A Dissertation Entitled Additive Manufacturing

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A Dissertation

entitled

Additive Manufacturing towards the Realization of Porous and Stiffness-tailored NiTi

Implants

Jason M. Walker

Submitted to the Graduate Faculty as partial fulfillment of the requirements for the

Doctor of Philosophy Degree in Biomedical Engineering

______Dr. Mohammad Elahinia, Committee Chair

______Dr. Sonny Ariss, Committee Member

______Dr. Sarit Bhaduri, Committee Member

______Dr. Christopher Cooper, Committee Member

______Dr. Matthew Franchetti, Committee Member

______Dr. Patricia R. Komuniecki, Dean College of Graduate Studies

The University of Toledo

May 2014

This document is copyrighted material. Under copyright law, no parts of this document may be reproduced without the expressed permission of the author. An Abstract of

Additive Manufacturing towards the Realization of Porous and Stiffness-tailored NiTi Implants

Jason M. Walker

Submitted to the Graduate Faculty as partial fulfillment of the requirements for the Doctor of Philosophy Degree in Biomedical Engineering

The University of Toledo

May 2014

This research is focused on the development of an additive manufacturing process for porous and stiffness-tailored nitinol implants. Selective laser melting (SLM) is an emerging additive manufacturing (AM) technology, which makes possible the production of 3D parts directly from metallic powders. AM enables production of 3D geometries that are not possible using traditional techniques. Other non-traditional manufacturing methods do not allow for precise control of pore geometry and distribution. With SLM, features such as engineered porosity, hollow parts, curved holes and filigree structures are suddenly realizable. Of particular interest in this research is the ability to engineer and control porosity.

The first stage of this study focused on understanding the processing parameters for manufacturing dense nitinol parts with a Phenix Systems PXM. A parametric analysis of process parameters on the quality and functionality of the SLM nitinol parts was conducted. First, a single track analysis was performed to understand the effect of basic processing parameters on the melting and re-solidification of powder. Next, a parametric

iii analysis of SLM process parameters, including laser power, scan velocity, and hatching spacing, on part density was carried out. Finally, a chemical analysis on impurity pickup during laser processing was performed. Based on the results of these studies, an optimal setup for processing nitinol on the Phenix Systems PXM was determined. Using the optimal parameter setup, SLM nitinol parts were manufactured for thermal, mechanical, and functional analysis. It was found that the transformation temperatures of SLM nitinol parts were approximately 10 C higher than the powder. This effect is attributed to nickel evaporation during laser processing. Mechanical properties were assessed in compression testing and determined to be very similar to properties of conventionally processed nitinol. The shape memory effect was also demonstrated, indicating that SLM nitinol retains its functional characteristics.

To improve the outcome of long-term metallic implant use, the mechanical properties of implants need to better match those of bone. Nitinol structures with engineered porosity were designed and manufactured by SLM to meet these requirements. It is shown that pore size, shape, and distribution can be customized by

CAD and manufactured by SLM. In this study, porous nitinol structures were manufactured with three different porosities (32%, 45%, and 58%) and featured interconnecting pores ranging from 0.6 mm to 1.0 mm in size. A reduction in elastic modulus based on porosity is demonstrated.

Acknowledgements

First, I would like to thank my advisor, Dr. Mohammad Elahinia, for the opportunity to work on this great project. I would also like to thank him for his guidance and support in all aspects of my research.

I owe many thanks to my committee members for taking time out of their busy schedules to provide the advice and direction required to complete this research. This work would not be possible without them.

I want to thank all of my colleagues and fellow researchers who I had the pleasure of working with in the Dynamic and Smart Systems Laboratory.

Finally, I want to thank my family. Mom and Dad – You inspire me every day.

Table of Contents

Abstract ...... iii

Acknowledgements ...... v

Table of Contents ...... vi

List of Tables ...... x

List of Figures ...... xi

1 Introduction…...... 1

1.1 Approach...... 2

1.2 Contribution...... 4

2 Metallic Implants ...... 5

2.1 Identification of Material Requirements ...... 6

2.2 Skeletal Architecture ...... 7

2.3 Effect of Surgical Implants on Adaptive Bone Remodeling ...... 9

2.4 Biocompatibility of Nitinol ...... 11

2.4.1 Biocompatibility of Porous Nitinol ...... 15

2.5 History of Nitinol in Medical Devices ...... 17

2.6 Nitinol in Orthopedic Applications ...... 18

3 Shape Memory Alloys ...... 22

3.1 History of Nitinol ...... 24

3.2 Transformation Behavior ...... 25

3.2.1 Phase Diagram ...... 26

3.3 Deformation Mechanisms ...... 32

3.3.1 Mechanical Behavior: Shape Memory and Superelasticity ...... 34

3.4 Physical Metallurgy ...... 38

3.5 R-Phase Transformations ...... 43

3.6 Manufacturing Limitations ...... 44

4 Additive Manufacturing ...... 46

4.1 History of AM ...... 47

4.2 Selective Laser Melting ...... 48

4.2.1 Background ...... 48

4.2.2 Phenix Systems PXM ...... 50

4.3 Existing Research on AM of Nitinol ...... 53

4.4 Process Parameters...... 53

4.4.1 Laser Power ...... 53

4.4.2 Beam Radius ...... 54

4.4.3 Hatch Spacing ...... 54

4.4.4 Layer Thickness ...... 54

4.4.5 Scan Strategy ...... 55

5 Parameter Setup, Part 1: Powder and Single Tracks...... 57

5.1 Powder Selection ...... 57

5.1.1 Composition ...... 57

5.1.2 Atomization...... 58

5.1.3 Characterization ...... 59

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5.2 Quality and Width of Single Tracks ...... 64

5.3 Mathematical Prediction of Single Track Width ...... 68

5.4 Summary...... 71

6 Parameter Setup, Part 2: Dense Parts ...... 72

6.1 Material Density...... 72

6.2 Chemical Analysis ...... 83

6.3 Summary...... 85

7 Thermal, Mechanical, and Functional Properties of SLM Nitinol ...... 89

7.1 Transformation Characteristics ...... 89

7.2 Compression Testing ...... 92

7.3 Cyclic Testing ...... 95

7.4 Summary...... 100

8 Designed Porosity in Nitinol Structures ...... 101

8.1 Review of Metallic Implant Requirements ...... 101

8.2 Computer-aided Design of Porosity...... 102

8.3 Compression Testing ...... 104

8.4 Summary...... 109

9 Conclusions and Recommendations for Future Work ...... 111

9.1 Specific Conclusions ...... 112

9.2 Recommendations for Future Work...... 114

10 Business Plan...... 116

10.1 Executive Summary ...... 116

10.2 Opportunity Rationale ...... 118

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10.3 The Company ...... 121

10.4 The Product ...... 122

10.5 Industry Overview ...... 124

10.6 Market Analysis ...... 128

10.6.1 Target Customers ...... 128

10.6.2 Market Size ...... 129

10.6.3 Competition...... 130

10.6.4 Estimated Market Share and Sales ...... 130

10.7 Marketing Plan ...... 131

10.7.1 Market Entry and Growth Strategy ...... 131

10.7.2 Price ...... 132

10.7.3 Sales Tactics and Service ...... 133

10.7.4 Distribution ...... 133

10.8 Product Development Plan ...... 134

10.8.1 Intellectual Property ...... 134

10.8.2 Product Design ...... 134

10.9 Manufacturing and Operations Plan ...... 136

10.10 Management Team...... 137

10.11 Overall Project Schedule...... 137

10.12 Critical Risks and Assumptions ...... 138

10.13 Financial Plan...... 140

References ...... 143

List of Tables

5.1 Transformation temperatures of nitinol powder ...... 62

5.2 Chemical impurity contents of nitinol powder ...... 63

5.3 Amplitude and exponent of the fitting functions by laser power...... 68

6.1 Impurity contents measured in powder and SLM nitinol part ...... 84

6.2 Optimal parameters for SLM processing of nitinol on a Phenix PXM...... 86

7.1 Transformation temperatures of SLM nitinol part and powder ...... 91

8.1 Dimensions of unit cells for porous structures and resulting pore formation ...... 103

8.2 Elastic modulus for typical implant materials and porous nitinol ...... 108

10.1 US 2004 knee implant market share, by company ...... 130

10.2 Financial calculations for Ortho3d ...... 142

List of Figures

3-1 Phase diagram of the transformation temperatures of nitinol ...... 27

3-2 Temperature-induced forward transformation in a zero stress state ...... 28

3-3 Temperature-induced reverse phase transformation in a zero stress state ...... 28

3-4 Twinned martensite is detwinned by an applied load ...... 29

3-5 Martensite remains in detwinned state when the applied load is removed ...... 30

3-6 Upon heating past Af, detwinned martensite transforms into austenite ...... 30

3-7 Superelastic effect ...... 31

3-8 (a) A crystal lattice and (b) resulting plastic deformation by an ordinary slip ...... 32

3-9 (a) A twinned crystal lattice and (b) resulting pseudoplastic deformation ...... 33

3-10 Stress-strain behavior of nitinol at different operating conditions ...... 37

3-11 Phase diagram of a Ti-Ni alloy ...... 39

3-12 Ms as a function of nickel content in quenched NiTi alloys...... 40

3-13 Effect of aging temperature on R-phase transformation start temperature ...... 43

4-1 Phenix Systems PXM Selective Laser Melting Machine ...... 49

4-2 Sequence of operations of the SLM process ...... 51

4-3 Visual representation of the Staircase Effect and support structures ...... 52

4-4 Basic scan strategy showing laser trajectories all in the same direction...... 55

4-5 Alternating x/y scan strategy ...... 55

4-6 Alternating x/y scan strategy with 90 degree rotation per layer ...... 56

5-1 Electrode Induction-melting Gas Atomization (EIGA) ...... 59

5-2 SEM micrograph of the nitinol powder ...... 60

5-3 Transformation temperatures of the nitinol powder ...... 62

5-4 Measured single track width as a function of scan velocity ...... 65

5-5 Micrographs of SLM single track ...... 67

5-6 Amplitude, a, of the fitting function as a function of laser power ...... 69

5-7 Calculated single track width as a function of scan velocity ...... 70

6-1 Cuboid blocks produced by SLM for density analysis ...... 74

6-2 Poor surface quality caused by too high energy input ...... 75

6-3 Relative density of SLM nitinol parts as a function of scan velocity ...... 75

6-4 Relative density of SLM nitinol parts as a function of energy density ...... 76

6-5 SLM nitinol samples mounted and polished for microscopy ...... 77

6-6 Comparison of pore formation in low and high density parts ...... 79

6-7 Optical micrograph of SLM nitinol ...... 80

6-8 Comparison of pore formation in two SLM nitinol parts ...... 81

6-9 Effect of hatch spacing on relative density of SLM nitinol parts ...... 82

6-10 Oxygen content in SLM nitinol parts...... 85

6-11 Micrograph of SLM nitinol single track with optimal parameter setup ...... 87

6-12 Micrograph of surface of SLM nitinol part with optimal parameter setup ...... 88

7-1 Transformation temperatures of SLM nitinol part and nitinol powder ...... 90

7-2 Stress-strain curve for SLM and conventional nitinol in the martensitic state ...... 93

7-3 Fracture of SLM nitinol samples at 45º with respect to the loading direction...... 95

7-4 Stress-strain curves of shape memory cycling to 300 MPa ...... 97

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7-5 Stress-strain curves of shape memory cycling to 900 MPa ...... 98

7-6 Cumulative irreversible strain in SLM nitinol during cyclic testing ...... 99

8-1 (a) Unit cell of porous design and (b) representative structure ...... 103

8-2 Porous SLM nitinol structures ...... 104

8-3 Micrographs of porous SLM nitinol structures ...... 105

8-4 Stress-strain curves to 12% strain for porous SLM nitinol ...... 106

8-5 Stress-strain curves to 2.5% strain for porous SLM nitinol ...... 107

8-6 (a) Stress-strain curve of SLM nitinol with 58% porosity and (b) fracture ...... 109

10-1 (a) Typical total knee replacement and (b) porous nitinol structure ...... 123

10-2 Estimated number of total knee replacements performed in the US ...... 126

10-3 Projected primary total knee arthoplasties in the US through 2030 ...... 127

10-4 Projected total knee revisions in the US through 2030 [206] ...... 127

10-5 Break-even analysis for Ortho3d ...... 141

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Chapter 1

Introduction

The aim of this research is to investigate the use of laser-based additive manufacturing to produce porous and stiffness-tailored nitinol implants. To improve the outcome of long-term metallic implant use, the mechanical properties of implant materials need to better match those of bone.

Orthopedic implants are used to restore and maintain proper biomechanical function in the body, yet current technologies utilize materials that possess drastically different mechanical properties than the biological materials they replace. Furthermore, these implants are introduced into highly unique environments (i.e. different patients), yet are mass designed with little regard for the individuals’ needs.

Most metallic implants are made out of 316 stainless steel, titanium (Ti-6Al-4V), or cobalt-chrome. Each of these materials can be 5 – 15 times stiffer than the surrounding bone tissue. During long-term fixation, this severely disrupts the physiomechanical processes which regulate bone remodeling. As a result, surrounding bone tissue degrades which leads to implant loosening and, ultimately, catastrophic failure.

To meet this unmet need, nitinol is identified as a more suitable material for use in long-term metallic implants. Nitinol features high strength, low stiffness, high recoverable strain, and good biocompatibility. Additionally, an additive manufacturing processing called selective laser melting (SLM) is identified as a better processing route than traditional methods. SLM circumvents current issues with manufacturing nitinol components (i.e. conventional machining of nitinol is incredibly difficult) and opens up possibilities for patient-specific implant design.

During the SLM process, a laser beam is used to manufacture a solid part by selectively binding powder particles through localized heating. The part is produced additively by manufacturing one cross-sectional layer at a time, which is digitally created by slicing a CAD model. Direct-from-CAD manufacturing allows for the design of structures with engineered porosity. In this way, the stiffness of porous structures can be modulated through the customization of pore size, shape, and distribution.

In summary, the primary objective of this research is to manufacture nitinol components with engineered porosity by selective laser melting to achieve desired mechanical properties.

1.1 Approach

SLM is a layer-based manufacturing process which utilizes direct melting of parts from a metallic powder-bed using a high-power fiber laser. The process features a considerable number of parameters which require optimization including powder characteristics, laser parameters, and layering technique. Even the smallest compositional

2 variances and all types of microstructural defects can strongly affect the integrity and functionality of nitinol parts.

Considerable work has been done to optimize the SLM process for materials such as titanium and stainless steel, but it is far from being established as a processing method for nitinol.

To this end, research is conducted on an SLM machine (Phenix Systems PXM) to develop and optimize the manufacturing method for dense nitinol components. A variety of analytical methods were used to evaluate and characterize fabricated AM nitinol components. These included microscopy, density analyses, differential scanning calorimetry (DSC), chemical analyses, and mechanical testing.

Once the optimal parameters were determined for the fabrication of dense nitinol components, the focus shifted to the design and fabrication of parts with engineered porosity. Nitinol components with various levels of CAD-designed porosity were manufactured and tested.

It is shown that we are able to fabricate functional nitinol parts with variable porosity and desirable mechanical properties.

1.2 Contribution

The contribution of this research is an understanding of the process parameters which affect the selective laser melting process for nitinol. Furthermore, it is shown that

SLM nitinol parts can be fabricated with CAD-designed porosity which can be used in medical implant technology to modulate the stiffness and facilitate bone ingrowth.

Chapter 2

Metallic Implants

In this chapter, the requirements, current state-of-the-art, and limitations of long- term metallic implant technology are reviewed. The field of biomaterials is constantly evolving and ever-expanding in a search for better solutions. Biomaterials are used in many different parts of the human body to repair and restore damaged organs. A non- exhaustive list of applications includes artificial heart valves, blood vessel stents, and joint replacements in the shoulders, knees, hips, and elbows [1].

Materials for orthopedic applications garner special attention due to the large number of injuries, demand for higher quality of life, and ageing world population [3].

Additionally, most orthopedic implants are long-term solutions. A vast array of materials including metals, alloys, ceramics, polymers, and composites have been developed to meet these needs. Unfortunately, there are still common applications which suffer from inadequate materials.

Most long-term metallic implants are made out of 316 stainless steel, titanium

(Ti-6Al-4V), or cobalt-chrome. Each of these materials can be 5 – 15 times stiffer than the surrounded bone tissue. During long-term fixation, this severely disrupts the physiomechanical processes which regulate bone remodeling. The presence of the

5 implant results in a redistribution of loads from the surrounding tissue to the implant itself. As a result, surrounding bone tissue degrades. Eventually, this leads to catastrophic implant failure.

2.1 Identification of Material Requirements

For a prosthesis to perform well inside the body, the implant materials must possess a wide array of properties. The required properties for the metallic component of a long-term implant include high strength, low elastic modulus, ductility, high corrosion and wear resistance, and good biocompatibility and ossseointegration [7].

High strength is needed to avoid fracture of the component. Ductility is required to avoid brittle failure. The material which is replacing bone should have an elastic modulus similar to that of bone. Large differences between the elastic moduli of implant material and bone can lead to stress-shielding, a primary cause of bone resorption and aseptic loosening. If an implant fails due to a mismatch in mechanical properties, then the material is considered biomechanically incompatible.

High wear and corrosion resistance are desirable to minimize the release of non- compatible metal ions by the implant into the body. These ions can accumulate in tissues near the implant or be transported to other parts of the body [9]. In either case, the release of ions has been shown to cause allergic and toxic reactions. Low wear resistance can also lead to aseptic loosening of the component. Finally, wear debris can cause an adverse inflammatory response in the tissue in which it is deposited.

Biocompatibility is required to ensure that the material does not induce any inflammatory or allergic reactions in the body. The two main factors that influence the biocompatibility of a material are the host response to the material and the materials degradation in the body. Highly biocompatible materials help optimize the rate and quality of bone apposition to the implant, minimize corrosion and the release of metal ions, minimize wear and the release of debris, and minimize disturbance to the biomechanical homeostasis in the bone [10].

Osseointegration refers to the ability of an implant surface to integrate with adjacent bone and to other tissues. It is believed that micromotion is a necessary mechanism for osseointegration. The inability of an implant to properly osseointegrate results in implant loosening. If the implant is not well integrated with the bone, an unwanted fibrous tissue develops between the bone and the implant.

In the following sections, it will be shown that nitinol possess a combination of high strength and ductility, low elastic modulus (comparable to bone), high corrosion and wear resistance, biocompatibility and ability to osseointegrate.

2.2 Skeletal Architecture

Bone tissue exists in a state of perpetual reconstruction. In response to mechanical forces, osteoclasts resorb bone and osteoblasts form new bone. This physiomechanical process, through which bone gradually changes its external geometry and internal structure, is called bone modeling [11]. Remodeling, or bone turnover, refers to resorption and subsequent formation of bone in packets [12]. During this process, the

7 activity of osteoclasts and osteoblasts is highly coupled and very little shape change occurs.

Bone remodeling is a lifelong process and progresses very gradually. However, during growth and in conditions of changed mechanical loading, bone modeling occurs.

During bone modeling, bone resorption and formation are not tightly coupled. New bone forms at a location different from the site of bone resportion and the shape of the bone changes [13]. A distinction must also be made between internal and external remodeling.

Internal remodeling is the adaptation of the density of bone tissue, characterized by a change of porosity [14]. External, or surface, remodeling refers to the apposition or removal of bone on the subperiosteal surface [15]. Bone remodeling serves many purposes. It constantly adjusts the bone architecture to meet mechanical needs and maintain skeletal strength homeostasis [16]. It also helps to repair microdamages in the bone matrix and prevent the accumulation of old bone [17].

Mechanical loading is the major functional influence on bone architecture [18].

The mechanostat hypothesis suggests that bone growth and bone loss are directly attributable to local strains in the bone, which are a function of mechanical loading [19,

20]. In response to relatively consistent loading, bone remodeling maintains a state of homeostatic balance. However, if the loading is abnormal, local osteoclasts or osteoblasts are activated. The resulting external modeling and internal modeling affect the geometry and density of the bone, respectively. This process is called strain adaptive bone modeling and continues until the mechanical stimulus is normalized [21].

The minimum effective strain (MES) constitutes a threshold that separates strains based on their ability to affect architectural bone remodeling [22]. Strains that exceed the threshold (50-500 microstrain) influence modeling and remodeling activities in ways that change the size and configuration of the bone [23]. When strains remain below the threshold, the bone enters into a state of “disuse-mode” remodeling, which is the body’s way of removing mechanically unnecessary bone. Disuse-mode remodeling increases local bone loss and decreases local bone strength [24].

A team of cells responsible for controlling bone remodeling and adaptation is called a basic multicellular unit (BMU) [25]. A BMU utilizes osteoclasts and osteoblasts in sequence to replace packets of old bone with new lamellar bone. BMU’s also detect and repair microscopic fatigue damage or microdamage (MDx), which is caused by repeated strains during normal loading [26]. However, when the strains exceed the MDx threshold (~3,000 microstrain), BMU’s are no longer capable of adequately repairing the bone. This results in a cycle of bone loss, increased strain, and generation of more microdamage. Sustained over an extended period of time, this gradual bone loss would eventually lead to fracture [26]. In young adults, the fracture strain of healthy bone is approximately 25,000 microstrain [28].

2.3 Effect of Surgical Implants on Adaptive Bone Remodeling

Bone architecture is also influenced by the presence of orthopedic implants. The performance of biomaterials is characterized by the host response to the implant and the ability of the implant to survive in the body [29]. The host response is characterized by three processes: Osteoinduction, osteoconduction and osseointegration. Osteoinduction is

9 the process by which bone formation is induced [30]. In a bone healing situation such as a fracture, the majority of bone healing is dependent on osteoinduction.

Osteoconduction refers to the ability of a material, such as an implant, to serve as a scaffold for bone attachment and growth. Materials with low biocompatibility, such as copper, silver and bone cement, exhibit little to no osteoconduction. Furthermore, an osteoconductive substrate must be porous to allow for the ingrowth of bone and fibrovascular tissue during osteoinduction [31-32].

Osseointegration, refers to a stable anchorage between bone and implant that results in a structural and functional connection [30]. For complete osseointegration, the bone must fill any gaps between itself and the implant, as well as repair any damage incurred during preparation of the implant site [33]. The primary difference between osseointegration and primary fracture repair is that, during the former, unification involves a foreign body [31].

Another aspect of bone remodeling around implants which must be addressed is the redistribution of mechanical loading. Conventional metallic implants are much stiffer than bone and, as a result, absorb a disproportionate amount of stress. This removal of normal stress from the bone by an implant is called stress-shielding [34]. This abnormal reduction of loading in the bone results in unwanted bone resorption and osteopenia.

2.4 Biocompatibility of Nitinol

The biocompatibility of a medical device refers to its compatibility with the human biological system. There are two primary factors that determine the biocompatibility of a material: the host reaction induced by the material and degradation of the materials in the body environment [35]. This section and the next will discuss the potential degradation of materials and the tissue response of the host. Titanium is highly biocompatible [36] and has been used in orthopedic implants since the 1950’s. A TiO2 film naturally forms on titanium surfaces providing the material with excellent corrosion resistance [37]. However, the biocompatibility of nitinol has been questioned due the potential leaching of nickel into the body [38].

Nickel is well known for its toxicity to humans and dissolution of Ni ions in the body could induce allergic, toxic and carcinogenic effects [39-41]. Epicutaneous tests have shown a dermatologic allergic reaction to nickel in about 5% of the general population [42]. In a rabbit study, pure nickel was implanted intramuscularly and into bone resulting in severe local tissue irritation, necrosis, and toxic reactions [43]. Due to the high Ni content in nitinol, it is theoretically possible that Ni-ions could dissolve out of the material. The corrosion resistance of nitinol in vivo controls the release of Ni-ions.

Corrosion resistance in metal alloys arises from passivation, a protective oxide coating that forms in contact with oxygen. It has been shown that after proper passivation, a 3-4 nm thick TiO2 layer forms on the nitinol surface [44]. Typically, the surface consists mainly of the TiO2 coating. In some cases, small amounts of nickel oxides (NiO and Ni-

2O3) and metallic Ni are also present. An In vivo study of nitinol clamps [44] showed that

11 only traces of metallic Ni were present on the surface after proper passivation. The samples were etched in a solution of HF-HNO3-H2O (1:2:3 vol%, 30 mins), pre- deformed, ultrasonically cleaned in ethanol, and sterilized by X-ray and room temperature to achieve proper passivation. No major change was observed in surface composition 4 and 12 months after implantation.

The first study on the biocompatibility of nitinol was done in 1973 by Cutright

[45]. Ni55Ti (55-nitinol) wire sutures were implanted subcutaneously in rats (n = 45). The tissue reaction to the implant was followed closely for 9 weeks. By the 6th week post-op, dense, relatively avascular fibrous connective tissue formed around the implant and tissue reaction was reported as minimal for the entire duration. It was concluded that the biocompatibility of 55-nitinol is indistinguishable from stainless steel and deep tissue implantation is feasible.

In 1976, high-purity nitinol femoral bone plates were implanted in dogs [46]. The plates were removed and examined after 3, 6, 12, and 17 months. In each case, there was no evidence of localized or generalized corrosion on the implants. Gross clinical, radiological, and morphological observations of the relevant tissue sites showed no adverse tissue reactions. The nitinol implant performance compared favorably the control group, implanted with cobalt-chromium, showing no discernible differences.

More recently, the biocompatibility of nitinol screws was compared to screws made from vitallium, titanium, duplex austenitic-ferritic stainless steel (SAF), and 316L stainless steel (1996) [47]. The screws were implanted in rabbit tibia for 3, 6, and 12 week periods. The biocompatibility results of nitinol compared well with the traditional

12 materials. However, the authors did observe diminished activity of bone cells and osteonectin synthesis in accordance with slower bone-remodeling around the nitinol screw.

Takeshita et al. [48] analyzed the bone reaction to cylindrical nitinol implants in the medullary canal of rat tibiae using a histomorphometric analysis (1997). The biocompatibility of the nitinol implant was compared to geometrically identical implants composed of titanium, anodic oxidized Ti (AO-Ti), Ti-6Al-4V, and nickel. The NiTi, Ti,

AO-Ti and Ti-6Al-4V implants were all progressively encapsulated with connective tissues over the course of the experiment (168 days). Histomertric analysis revealed no significant differences among tissue reactions between Ti, AO-Ti and Ti-6Al-4V.

However, the NiTi implants show significantly lower percentage bone contanct and bone contact area than the other Ti-alloys. Furthermore, the Ni implants showed no bone contact at any time during the experimental period.

In 1999, Ryhänen et al. [49] performed a study on the effects of NiTi on osteotomy healing, bone mineralization, and the remodeling response. NiTi and stainless steel intramedullary nails were used to fix femoral osteotomies in 40 rats. Bone healing was examined with radiographs, peripheral quantitative computed tomography (pQCT), and histology. The corrosion of the used implants was analyzed by electron microscopy

(FESEM). Analyses of the implants were performed after 2, 4, 8, 12, 26, and 60 weeks.

In the early stages (4 and 8 weeks), there were more healed bone unions in the NiTi group than in the stainless steel group. At the later stages, there was no noticeable difference in bone healing between the two groups. Both the NiTi and stainless steel groups achieved 100% healed bone unions by the week 60. It was concluded that NiTi

13 does not impair the osteotomy healing response, consolidation, mineralization, or remodeling of bone compared to stainless steel. Furthermore, harmful Ni-ion release due to corrosion of NiTi intramedullary nails is not likely.

In another study by Ryhänen in 1999 [50], new bone formation, modeling and cell-material interface responses induced by NiTi periosteal implants were evaluated.

Periosteum is a fibrous membrane that lines the outer surface of bones. Here, the femoral periosteum was contacted with implants made of NiTi, Ti-6Al-4V, and stainless steel.

Bone formation and resorption parameters were determined using histomorphometry with digital image analysis. Bone formation at 12 and 26 weeks showed no statistical difference between NiTi and the other materials. The histological response of the soft tissues around the NiTi implant was clearly non-toxic and non-irritating. The authors concluded that the NiTi had no negative effect on total new bone formation.

An ecoptic bone formation assay was used to study effects of NiTi implants in

2000 [51]. NiTi, stainless steel, and Ti-6Al-4V materials were implanted for 8 weeks under the fascia of the latissimus dorsi muscle in 3-month old rats. It was found that total bone mineral density values were nearly equal between the control and the NiTi group, while the stainless steel and Ti-6Al-4V groups showed lower bone mineral density.

Furthermore, the NiTi group showed the largest proportional cartilage and new bone area.

2.4.1 Biocompatibility of Porous Nitinol

The biocompatibility of monolithic nitinol implants is well established. However, due to the different manufacturing techniques and increased surface area, there could be an increased risk of Ni leaching in porous nitinol implants. To this end, many studies have been conducted to investigate potential cytotoxicity of porous nitinol [52-67]. In

1995, Simske [52] conducted a study on uncoated porous nitinol craniofacial implants.

Rabbit subjects were evaluated 2, 6, and 12 weeks after implantation. It was observed that soft and connective tissues adhered to the implants before the 2 week evaluation and there was no indication of tissue inflammation. Microhardness and histologic assessments indicated new bone growth into the implants had similar properties to the surrounding bone [52]. A similar study conducted in 1999 [53] found similar results in tibial osteotomies. No adverse effects were observed and bone tissue demonstrated good healing of the osteotomy.

The surface corrosion and nickel release resistance of porous nitinol implants was evaluated in 2002 by Assad [56]. Porous nitinol interbody fusion devices (Actipore™;

Biorthex Inc) were compared to a conventional nonporous Ti-6Al-4V cage. The devices were implanted into sheep and compared after 3, 6, and 12 month periods. Using scanning election microscopy (SEM), no evidence of surface corrosion was found in any of the implants. Blood nickel levels were monitored and found to be within acceptable levels at all times. Furthermore, there was no difference in nickel content in implant- adjacent tissue and detoxification organs between the subjects with nitinol implants and the controls [56]. Histomorpometric and radiological analyses were also performed. The

15 authors found that radiological fusion scores were consistently superior for the nitinol implants [57]. They conclude that porous nitinol constitutes an excellent substrate for osteogenic cell integration and represents a new osteoconductive biomaterial with better fusion characteristics than conventional titanium cages.

Sufficient bone contact and ingrowth is equally important as minimizing allergic, toxic and carcinogenic effects when characterizing the biocompatibility of a bone implant. Without the former, the latter is irrelevant. Osseointegration of porous nitinol was evaluated [58] during a 12-month study on intervertebral lumbar implants in sheep.

Radiological analysis showed that the majority of porous nitinol implants (15/16,

93.75%) demonstrated bone integration and apposition, which was much better than the titanium control group (4/16, 25.0%). Furthermore, bone density surrounding the nitinol implants indicated adequate bone remodeling. Bone growth into porous metal surfaces depends on many factors [59, 60], of which one is the porosity of the surface. Kujala et al. studied the effect of porosity on osseointegration [61] of nitinol implants in rat bone.

They found that implants with a porosity of 66.1% and mean pore size (MPS) ± 30 μm showed the best bone contact compared to 59.2% (MPS 272 ± 17 μm) and 46.6% (MPS

505 ± 136 μm) porosities. However, the 46.6% porosity was not significantly inferior in bone contact.

2.5 History of Nitinol in Medical Devices

One of the unique benefits of the human body, as a medium for nitinol applications, is that it offers an isothermal environment. Early in its development, nitinol was seen as a potential material for biomedical applications. However, throughout the

1970’s and 1980’s, there were many obstacles which blocked rapid development of nitinol applications [68]. The non-linear tensile properties, hysteresis, fatigue, and adiabatic heating and cooling effects were not yet fully understood [69]. As a result, most of the large companies at the time lost interest in nitinol and refocused on traditional materials. Nevertheless, the scientific community remained steadfast and in the 1990’s, a variety of circumstances converged to stimulate the emergence of nitinol in the medical industry [68]. There was a large push in the industry for device development which would allow less-invasive procedures, manufacturing methods such as microtubing and laser cutting were available, and material scientists and companies were negotiating the

‘release’ of technology to product designers and doctors. This culminated in the first commercially successful nitinol medical device with self-expanding stents. More importantly, it shifted the industry’s perception of nitinol back into favorable light.

The earliest nitinol devices in the medical industry were catheter-deployed. They took advantage of the enormous elasticity of nitinol to accomplish jobs that regular materials cannot do. The first such commercial product was the Homer Mammalok needle/wire localizer [24,70]. The device used a superelastic hook which could be withdrawn into a straight catheter and reform its hook shape after deployment. Later, devices incorporating thermal deployment were developed. One such device is the Simon vena cava filter [71-72]. The device is inserted into the vena cava in a straight-wire

17 configuration and transforms its shape through the shape memory effect when it equilibrates to body temperature. In its transformed shape, it acts as a filter to catch embolized blood clots in the vena cava. Since its inception into the medical world, nitinol devices have become commonplace in the orthodontic [73-76] and vascular fields [77-

80]. More recently, researchers have focused on developing orthopedic applications for nitinol which utilize its unique mechanical properties [81].

2.6 Nitinol in Orthopedic Applications

Shape memory compression staples [82] were developed for the treatment of fractures. In this application, the staple is deformed in the low temperature martensite phase prior to implantation. After implantation, the temperature of the body induces the shape memory effect. Because the recovery is physically constrained, the staple provides a sustained compression force on the fracture line, facilitating repair. In one study [82], patients with intra-articular fractures were treated, including fractures of the patella, malleoli, olecranon, lateral condyle and capitellum of the humus, and the tibial plateau.

Another study reported the successful fixation of a mandible fracture [83] using a nitinol staple.

Another application involves implanting nitinol rods for the correction of scoliosis

[84-89]. In 1975, Schmerling [84] proposed the use of a Harrington rod made of nitinol.

During the proposed surgery, the spine would be distracted by an external device and a nitinol rod would be inserted. The rod would be attached to the transverse processes of the vertebrae and adjusted to apply the correct force to the spine. In 1993, Sanders [85-

86] performed an in vivo study on goats with experimental scoliosis. Prior to

18 implantation, the rod was deformed to accommodate the existing curvature of the spine.

The implanted rods were then transformed post-operatively using radio frequency induction heating to apply corrective forces to the spine. It was observed that the scoliosis was corrected in all six subjects. Another experimental study was performed on pigs in

2002 by Wever [87]. In this study, the shape memory effect was induced perioperatively and was used to induce scoliosis. Macroscopic inspection after 3 and 6 months showed that the device was almost completely overgrown with new bone formation. No corrosion was observed and histologic examinations showed no evidence of a foreign body response.

One particular advantage of the shape memory effect may be in the development of minimally invasive surgical procedures. A spinal vertebral spacer has been proposed

[72] that uses the shape memory effect replace a damaged intervertebral disc. The implant would be deformed prior to surgery to facilitate its insertion between vertebrae during surgery. Once inserted, the proposed device would recover its original shape via the shape memory effect. In this application, there would be no external heat source; the shape memory effect is induced entirely by equilibration to body temperature.

Interspinous implants utilizing the superelastic effect have also proposed [91] and studied in simulation. Finite-element models show that interspinous nitinol implants have a more physiological stiffness, while interspinous titanium implants are excessively stiff and exhibit plastic strains under physiological loads.

Controlled bone modeling during implantation of intramedullary nitinol nails [92-

93] was studied in 2002. Pre-shaped intramedullary nitinol nails were implanted in the medullary cavity of femurs in rats. The straightened nails were implanted in the

19 martensitic phase and, upon warming to body temperature, transformed to the austenite phase which had a preset bent shape. After 12 weeks, expected bending of the bone, bone thickening and increase in cortical area were observed. A control group with straight nails showed significantly less increase in bone cross-sectional area and cortical thickness, indicating that the bending force of the functional nail was the primary cause.

One of the most recent orthopedic applications of shape memory alloys involves the development and use of porous foams. A method for manufacturing porous nitinol

[95] was first patented in 1972; however, extensive research was not conducted on the biocompatibility and corrosion resistance until recently [52-65]. The use of porous nitinol in orthopedic applications is promising for many reasons. Important properties of porous nitinol include: high strength – to prevent fracture; relatively low stiffness – to minimize stress shielding effects; high toughness – to avoid brittle failure; shape memory behavior

– to facilitate implant insertion and/or ensure good mechanical stability within the host and/or provide constant force; and good biocompatibility – to ensure healthy osseointegration [81].

Actipore™ is a porous nitinol foam produced by the company Biorthex Inc

(Montréal, Canada). Actipore™ is currently utilized in two implant systems [96-97]. The

Actipore™ PLF System, introduced in 2001, is an interbody fusion device for lumbar fusion proceduces. The Actipore™ ACF system, introduced in 2005, is a similar device for cervical fusions. Worldwide, there have been more than 1,000 implantations using

Actipore™ [98]. The devices are not licensed in the United States. The functionality of the Actipore™ interbody fusion implant was evaluated in sheep [58] and compared to a conventional Ti-6Al-4V intervertebral cage (2004). The osseointegration capacity was

20 assessed after 3, 6, and 12 months of implantation. Two-dimensional radiology, CT scans, and histological testing showed that the porous nitinol implant is superior to the

Ti-6Al-4V cage in terms of osseointegration.

Chapter 3

Shape Memory Alloys

Shape memory alloys (SMAs) are a class of smart materials that exhibit a direct coupling between temperature and mechanical response. Direct coupling implies that either stimulus, thermal or mechanical, can serve as an input and the other as an output.

The unique ability of active materials like SMAs to act as solid-state actuators and sensors has catapulted them to the forefront of scientific innovation [99].

Other active materials that exhibit a direct mechanical coupling include piezoelectrics and electrostrictives (coupled with electric fields), and peizomagnetics and magnetostrictives (coupled with magnetic fields). Of all active materials that exhibit a direct coupling, SMAs possess the highest actuation energy density [100]. Their extraordinary power density allows for the development of lightweight, high force/displacement sensors and actuators. Additionally, their inherent mechanical simplicity adds to their favorability over conventional systems such as pneumatics, hydraulics, and solenoids.

The thermoelastic behavior of shape memory alloys is the result of a diffusionless, solid-solid, reversible phase transformation between a high temperature

22 austenite phase and a low temperature martensite phase [101]. A critical transformation temperature defines the relative phase stabilities and is dependent on material composition. The results of this transformation are two phenomena that characterize all

SMAs: the shape memory effect and superelasticity. The shape memory effect (SME) occurs when an SMA is plastically deformed in its martensitic phase and then recovers its initial shape through a heating process. Superelasticity is observed when an SMA is deformed in its austenitic phase and immediately recovers its initial shape upon unloading. Some of the known shape memory alloys include Ag-Cd, Au-Cd, Cu-Al-Ni,

Cu-Sn, Cu-Zn-(X), In-Ti, Ni-Al, Ni-Ti, Fe-Pt, Mn-Cu, and Fe-Mn-Si.

Shape memory phenomena were first reported in Au-Cd systems by Ölander in

1932, who famously described its behavior as rubber-like [102-104]. His observation became the first record of superelasticity in metals and ultimately created a new field of study which would become shape memory alloys. Shortly thereafter, in 1938, Greninger and Mooradian reported the formation of a temperature dependent martensite phase in

Cu-Zn alloys [105]. The crystallographic and kinetic basis for the thermal reversibility of matertensites, which governs shape memory behavior, was reported by Kurdyumov and

Khandros in 1949 [106]. Further research was conducted on the behavior of martensitic twins in Au-Cd by Chang and Read [107] (1951) but it wasn’t until thermoelastic phenomena were discovered in Ni-Ti systems that shape memory alloys gained mass popularity.

3.1 History Of Nitinol

The discovery of thermoelastic behavior in nickel-titanium alloys was made quite accidentally in 1963 [108-111]. The fortuitous discovery was made by Buehler and his colleagues while investigating the heat and corrosion resistance of intermetallic materials for thermal shields at the Naval Ordnance Laboratory (NOL). While examining the thermal properties of different materials, it was observed that NiTi exhibited very unusual properties including temperature dependent acoustic damping characteristics. Struck by this bizarre discovery, the group maintained interest in the alloy. Still, it wasn’t until a routine laboratory meeting that an associate technical director curiously heated a bent corrugated strip of NiTi with his pipe lighter. To everyone’s astonishment, the NiTi straightened itself out [112]. Following this surprising observation, Buehler and his colleagues continued to study NiTi and eventually coined the terms shape memory and nitinol (Ni-Ti-NOL).

Nitinol quickly gained scientific popularity and its quest to become the most widely-utilized shape memory alloy was made with haste. A fundamental understanding of the crystallographic mechanisms that drive the SME [113-116], as well as stress- induced martensitic transformations [117-118], was completed by the 1970’s. Much of the work was done on Cu-based alloys due its low cost and ease of processing single crystals. However, Cu-based alloys are quite brittle in their polycrystalline state and nitinol possesses huge advantages in terms of functional properties [119]. Some of the advantageous properties of nitinol include high ductility, high strain recoverability, excellent corrosion resistance, stable transformation temperatures, relatively low elastic

24 anisotropy, high biocompatibility and good magnetic resonance imagining opacity [121-

122].

To date, the largest application of superelastic nitinol has been in the biomedical industry [121,123-128], with roles in orthopedic and cardiovascular surgeries, and orthodontics. Many eyeglass frames and toys also utilize superelastic nitinol [129]. Shape memory nitinol is widely utilized in actuators and sensors in micro-electronic mechanical systems (MEMS) [130], as well as in solid-state heat engines, shrink-fit pipe couplers for aircraft, and other applications [131].

3.2 Transformation Behavior

The mechanism that governs the shape memory effect and superelasticity in nitinol is a diffusionless, solid-solid, first-order phase transformation [132-141]. The first-order phase transition is defined by the presence of latent heat of the transformation that expresses the amount of heat required for phase transformation. There are two distinct, temperature-dependent solid phases in nitinol. The parent phase, austenite, is stable at high temperature and is an interpenetrating cubic B2 crystal structure (CsCl type). The low temperature product phase, martensite, is a low-symmetry monoclinic

B19’ structure. The forward transformation can occur as a single-step, B2 → B19’, or as a two-step transformation, B2 → R → B19’. The intermediate R-phase is a trigonal crystal structure [132-135]. Whether the R-phase occurs is dependent on composition and material processing. The reverse transformation, from martensite to austenite, is always a single step transformation; B19’ → B2 [132].

In addition, the martensite structure can organize itself in two ways: twinned and detwinned. Individual martensite crystals can have different orientation directions. These orientations are called variants. In nitinol, there are 24 possible martensitic variants

[135]. Twinned martensite consists of a combination of many self-accommodating martensitic variants. Detwinned martensite consists of a single, dominant variant.

3.2.1 Phase Diagram

In addition to being temperature dependent, the phase stability is also dependent on stress. Figure 3.1 is a visual representation of the transformation temperatures of nitinol as a function of applied stress. This diagram is called a stress-temperature phase diagram and shows the transformation zones for an SMA with fixed composition. For simplicity, we only consider a single-step transformation from B2 austenite to B19’ martensite. During the nitinol transformation, there are four characteristic transformation temperatures: austenite start (As), austenite finish (Af), martensite start (Ms), and martensite finish (Mf). The transformation temperatures increase in response to applied stress, indicated in the figure by diagonal lines. The transformation behavior of nitinol is hysteretic. That is, the temperature and stress states of the material are not adequate to determine the phase stability. The loading history must also be known.

Figure 3.1: Phase Diagram of the transformation temperatures of nitinol as a function of applied stress

The phase diagram can be used to illustrate the superelastic and shape memory behaviors of nitinol. Consider a specimen at a temperature above Af in a zero stress state.

This specimen is fully austenitic. Subsequent cooling causes the austenitic material to begin forward transformation to twinned martensite as it passes the Ms temperature.

When the temperature reaches Mf, the material is entirely transformed into twinned martensite (Figure 3.2). During this transformation, the martensite crystals form in a combination of twenty-four self-accommodating variants such that no macroscopic deformation occurs. Heating the material from this state induces reverse phase transformation from twinned martensite to austenite (Figure 3.3). The reverse transformation begins at As and concludes at Af. To illustrate the hysteretic behavior, consider a specimen at zero stress at a temperature, T, between Ms and As. During

27 forward transformation, the specimen is still fully austenitic. At the same temperature during reverse transformation, the specimen is fully martensitic.

Figure 3.2: Temperature-induced forward phase transformation in a zero stress state (after [100])

Figure 3.3: Temperature-induced reverse phase transformation in a zero stress state (after [100])

The shape memory effect (SME) involves steps in which the material is physically deformed and then recovers its original shape via heating. Consider a specimen that begins in its parent shape, which is fully austenitic. The first

28 transformation is a zero-stress cooling stage from austenite to twinned martensite (Figure

3.2). Subsequent loading of the specimen forces the different martensite variants to align into a single, geometrically favorable variant (Figure 3.4). This process is called detwinning and a macroscopic deformation is observed. Detwinning begins and completes at critical stresses, σs and σf. Upon release of the applied load, the specimen remains in its detwinned state (Figure 3.5). Finally, the specimen is heated Af, inducing a reverse transformation to its parent phase, austenite (Figure 3.6). During this transformation, the macroscopic deformation is recoved and the specimen reverts to its original shape [100,132].

Figure 3.4: Twinned martensite is detwinned by an applied load; macroscopic deformation is observed (after [100])

Figure 3.5: Martensite remains in detwinned state when the applied load is removed; macroscopic deformation remains (after [100])

Figure 3.6: Upon heating past Af, detwinned martensite transforms into austenite; macroscopic deformation is recovered (after [100])

The superelastic effect (SE) is a stress-induced transformation that occurs at high temperature (Figure 3.7). Consider a specimen in its parent phase above Af. Isothermal loading induces a macroscopic strain which is accompanied by a phase transformation to detwinned martensite. Subsequent unloading immediately results in the recovery of macroscopic deformation and a transformation back to austenite.

Figure 3.7: Superelastic effect: Isothermal loading above Af induces a reversible phase transformation from austenite to detwinned martensite (after [100])

3.3 Deformation Mechanisms

In both shape memory and superelastic applications, axial strains of up to 8% are fully recoverable for nitinol [34]. Compared with typical metals that are commonly used in biomedical applications, this value is incredible. In response to a sufficiently large applied load, all crystalline materials will undergo a macroscopic deformation. The early stages of deformation involve the stretching of bonds between atoms. This type of elastic deformation is the only recoverable deformation in typical materials and is characterized by very small strains. Stainless steel, for example, possesses approximately 0.1% elastic strain recoverability. Beyond the elastic strain limit in typical materials, molecular bonds break and plastic deformation occurs by the movement of dislocations. The motion of a large number of dislocations is called slip. In response to the applied shear force, bonds break along a slip plane and part of the crystalline lattice is displaced with respect to another (Figure 3.8). The global crystal structure remains unchanged but a macroscopic and irreversible shape change has occurred.

Figure 3.8: (a) A crystal lattice and (b) resulting plastic deformation by an ordinary slip dislocation along a slip plane

In a shape memory alloy such as nitinol, the early stages of deformation are also characterized by an elastic regime in which the molecular bonds are stretched. However, beyond the elastic strain limit a pseudoplastic phenomenon is observed. The pre- deformed material consists of martensite twins: martensite variants which mirror each other across a twin plane. Instead of a slip step, it is more energetically favorable for one of the reflections to reorient so that it matches its twin (Figure 3.9). This deformation is called detwinning and is a pseudoplastic deformation because it is fully recoverable through heating [142-143].

In a superelastic case, stress-induced martensitic (SIM) transformation gives rise to the deformation upon loading [144]. However, in the absence of stress the martensite is unstable at this temperature. Therefore, the material automatically transforms back to austenite upon unloading.

Figure 3.9: (a) A twinned crystal lattice and (b) resulting pseudoplastic deformation by detwinning along a twin plane

3.3.1 Mechanical Behavior: Shape Memory and Superelasticity

The deformation mechanics in shape memory and superelastic nitinol manifest in interesting stress-strain behavior. Due to the thermoelastic and hysteretic nature of shape memory alloys, the mechanical behavior of nitinol is dependent on temperature and initial phase conditions. This behavior in response to isothermal loading is broken down into four categories (Figure 3.10):

(a) Martensite, T < As

When the austenite phase cools down below Mf in a stress-free state, the material undergoes a martensitic transformation without any macroscopic shape change that is achieved through self-accommodation of martensite variants (Figure 3.2). The martensite remains stable while the temperature stays below As. In the absence of stress, the material consists of many martensite variants in twins. Upon loading, the martensite initially deforms elastically. When the applied stress increases above a critical point, σs, the martensite twins begin detwinning into a single preferred variant (Figure 3.4). The preferred variant depends on the loading direction and crystal orientation of the texture.

The variant in the most energetically favorable orientation becomes the preferred variant.

During the detwinning process, the stress rises very slightly creating a plateau. At a critical stress level, σf, the detwinning process completes, the martensite consists of a single variant, and the stress plateau terminates. Further application of stress induces elastic deformation according the martensite modulus. The detwinned martensite is thermodynamically stable below As. Therefore, upon unloading of the applied stress,

34 reorientation of the martensite into a twinned state does not occur and only a small amount of the strain is recovered (Figure 3.5). The residual strain and initial macroscopic shape is recovered when the martensite is transformed to austenite by heating above Af

(Figure 3.6). Subsequent cooling below Mf in a stress-stree state will return the specimen to its twinned martensite phase.

(b) Austenite, Ms < T < As)

Consider a specimen that begins in the austenite phase at high temperature. The specimen is cooled to a temperature between Ms and As (Figure 3.1). In this temperature state, martensite transformation has not started and the specimen remains austenitic.

Upon isothermal loading, the austenite initially deforms elastically. Above a critical stress level, the austenite yields and starts transforming to martensite. The stress-induced transformation occurs at near-constant stress, forming a stress plateau. The plateau occurs due the martensite phase’s ability to form variants during the transformation. When transformation is complete, the martensite phase deforms elastically according the martensite modulus. In this temperature state, below As, the martensite is thermodynamically stable. Unloading the specimen results in the same behavior as case

(a).

Partial strain recovery is observed when the operating temperature is between As and Af (Figure 3.1). Here, the specimen has been cooled from a temperature above Af to ensure that it begins in a fully austenitic state. Loading results in elastic deformation according to the austenite modulus followed by a stress-plateau during stress-induced

35 transformation to martensite and finally elastic deformation according to the martensite modulus. The martensite is stable above As due to the applied stress. However, during the unloading process, the martensite becomes unstable and a reverse transformation to austenite occurs. This reverse transformation is accompanied by the development of another stress plateau. The reverse transformation stress plateau occurs at a much lower stress level than during forward transformation. In this temperature state, though, complete reverse transformation does not occur there is some residual strain. Again, recovery of the residual strain is possible by heating the material above Af.

(d) Austenite, T > Af

The final case is illustrates the superelastic effect and occurs above Af (Figure

3.7). The behavior is entirely similar to case (c) until the latter part of unloading. At temperatures above Af, complete reverse transformation to austenite occurs upon the removal of stress. At this time, the unloading stress plateau terminates and the remaining strain is recovered elastically according to the austenite modulus. Recoverable strains in superelastic nitinol have been observed as high as 11%, although in most conditions the maximum is found to be around 8%. It is important to be aware of the recoverable strain limits. Strain beyond a critical limit in the second elastic region will induce crystallographic slip, thus leading to unrecoverable plastic strains.

Figure 3.10: Stress-strain behavior of nitinol at different operating temperatures and initial phase.

(a) Martensite, T < As (b) Austenite, Ms < T < As (c) Austenite, As < T < Af (d) Austenite, T > Af

The stress-induced martensitic transformation that drives the superelastic effect is a shear-like mechanism. As shown in the previous section, it is possible to stress-induce martensite at temperatures above Ms. The Clausius-Clapeyron thermodynamic relationship summarizes the transformation [132,145]:

푑휎 훥푆 훥퐻∗ 3.1 = − = − 푑푇 휀 휀푇

Where σ is a uniaxial applied stress, ε is transformation strain, ΔS is entropy of transformation (per unit volume), and ΔH* is the enthalpy of transformation (per unit volume). During the forward transformation from austenite to martensite, ΔS and ΔH* are negative. Thus, for stress-induced martensitic transformation, the slop of dσ/dT must be positive. Furthermore, the temperature dependence of transformation stresses is given by:

푑휎 3.2 푀 = 퐶 푑푇 푀 푑휎 3.3 퐴 = 퐶 푑푇 퐴

Where CM (forward transformation) and CA (reverse transformation) are phase-dependent constants. Equations 3.2 and 3.3 indicate that the transformation temperatures during forward and reverse transformation are linearly affected by the application of stress.

3.4 Physical Metallurgy

The elemental composition of nitinol is not the only driving factor of phase transformation behavior. Thermomechanical treatments have a direct and substantial effect on the formation of dislocations and precipitates, which in turn influence the deformation behavior, shape memory effect, and the occurrence of R-phase. However, elemental composition is the most fundamental characteristic of the material. So, while it cannot necessarily be used as an accurate predictor of material behavior by itself, it is the most fundamental starting point. The effect of thermomechanical treatments will follow.

The Ni-Ti alloy system is quite complicated. Single phase “TiNi” is only present in alloys near equiatomic composition (Figure 3.11) [146]. The TiNi region is bounded by Ti2Ni to the left and TiNi3 to the right. Furthermore, the TiNi region is relatively vertical on the Ti-rich side right about 50-50 atomic percent. On the Ni-rich side, nickel solubility is as high as 56% at 1115 °C. From this point, the boundary on the Ni-rich side decreases with decreasing temperature. Therefore, nitinol does not have a fixed composition. The alloy cannot dissolve excess Ti, but it can dissolve up to 6% excess Ni at high temperatures. Quenching from high temperature can preserve the solid solution to room temperature without precipitation [132].

Figure 3.11: Phase diagram of a Ti-Ni alloy [146]

In quenched Ni-Ti alloys, the transformation is a single-step, B2 → B19’, and the transformation temperature is highly dependent on the composition. Change in nickel concentration by 0.1 at% in Ni-rich alloys can shift Ms by more than 10 °C [147]. Tang

[148] collected data from [149-155] and predicted the martensite transformation start temperature for various compositions. Figure 3.12 shows this data as a function of nickel content [148]. On the Ti-rich side, the transformation temperature is virtually constant at approximately 60 °C. However, on the Ni-rich side, the transformation temperature decreases drastically with increasing nickel content. This relationship is incredibly important in the development of nitinol components. While there are other motivating characteristics, the nickel content can be used to help tailor the transformation temperature of a component to a specific application.

Figure 3.12: Approximate martensite start temperature (Ms) as a function of nickel content in quenched NiTi alloys (predicted by Tang [148])

In most cases, precise control of the transformation temperature is required to guarantee the functional quality of the nitinol component. However, as shown in Figure

3.12, the transformation temperature is highly sensitive to composition. Therefore, composition is used as a general starting point and then thermomechanical treatments are used to fine tune the transformation temperature. It is important to note that the following methods do not apply to Ti-rich alloys.

The transformation temperature of quenched NiTi alloys can be affected by the introduction and manipulation of precipitates in the Ti-Ni matrix. Nickel-rich precipitates can be introduced into nickel-rich alloys via aging treatments by diffusional transformations. At lower temperatures and shorter times, Ti3Ni4 precipitates appear.

Under certain conditions, further aging at higher temperatures and longer times causes the Ti3Ni4 phase to be absorbed into the matrix while a Ti2Ni3 phase appears. Still further aging at even higher temperatures and longer times can cause the Ti2Ni3 phase to be absorbed into the matrix while a TiNi3 phase appears. Technically, the TiNi3 phase is the equilibrium phase while the former two are intermediate phases. However, under normal aging conditions, only the metastable Ti3Ni4 precipitate is observed and it is quite stable below 600 °C. The upper limit of the Ti3Ni4 phase has been shown to be around 680 °C

[34]. Therefore, manipulation of the Ti3Ni4 precipitate is used to tune the transformation temperature.

The technique for adjusting the transformation temperature of Ni-rich alloys was first reported by Horikawa in 1989 [156]. It was discovered that alternatively aging Ni- rich alloys between two aging temperatures resulted in the repetitive changing of transformation temperature. That is, aging at a high temperature (550 °C) decreased Ms.

Subsequent aging at a lower temperature (440 °C) increased Ms. This process could be used to adjust the transformation temperature up and down for many cycles. It was also determined that aging time had a negligible effect on the outcome. Therefore, the transformation temperature could be controlled by aging temperature alone, and with only a very slight dependence on number of cycles.

After extensive research, Zhang et al. [157] found that this phenomenon is due to a temperature-induced change in the B2 matrix composition. During the aging process, a precipitation reaction occurs, which results in the formation of Ti3Ni4 precipitates and a decrease in Ni content in the TiNi matrix. Due to the phase equilibrium between the TiNi matrix and Ti3Ni4 precipitates, aging at lower temperatures results in a larger decrease of

Ni content in the TiNi matrix [158]. Therefore, aging at lower temperatures results in increased transformation temperature and vice versa (Figure 3.13). While the data presented shows the effect on R-phase transformation start temperature (Rs), the effect on martensite transformation start temperature is analogous.

Figure 3.13: Effect of aging temperature on R-phase transformation start temperature (after [158])

3.5 R-Phase Transformations

In some cases, the transformation from B2 results in another martensite called R- phase. Initially, R-phase was considered a precursor to B19’ martensite [159-161], but we now know that the B2 → R-phase transformation is a distinct martensitic transformation that competes with the B2 → B19’ transformation. The martensite with the smaller transformation strain is the phase which will develop [132]. In the case that B2 transforms into R-phase, continued strain development will force the R-phase to transform into B19’ to further lower free energy. In this situation, the forward transformation becomes a two-stage transformation B2 → R → B19’. The R-phase is not

43 thermodynamically stable in the reverse direction. Therefore, the reverse transformation always occurs as a single step from B19’ → B2.

In a previous section, it was noted that in quenched NiTi alloys, the phase change is a single-step martensitic transformation from B2 to B19’. The R-phase only occurs when the presence of defects of secondary phases interferes with the B2 → B19’ transformation [162]. This can be achieved by:

a) Introduction of dislocations (through work-hardening)

b) Formation of precipitates in Ni-rich alloys (though age-hardening)

c) Ternary alloying elements such as Fe, Al, etc.

The recoverable strain during the B2 → R-phase transformation is only on the order of 0.5 - 1% [140,141]. However, the second-stage transformation from R-phase →

B19’ martensite is still capable of producing recoverable strains in excess of 6% [135].

The B2 → R transformation generates a small stress plateau on the stress-strain curve.

Compared to the R → B19’ stress plateau, though, this plateau is very brief and exists at a much lower stress/strain level. \

3.6 Manufacturing Limitations

Machining nitinol is incredibly difficult. Stress induced martensitic transformation, work hardening, spring back effects, burr formation and adhesion can quickly degrade the work piece quality and can cause considerable tool wear [163,164].

Also, the thermal conductivity of nitinol is very poor, making it difficult to remove heat.

Grinding, therefore, results in high heat generation, which leads to crack formation on the surface of the part. Laser cutting, which provides a tool-free method of fabricating precision parts, involves challenges such as creation of intermetallic phases (Ti2Ni), and crack formation in the heat affected zone. Laser cutting is also limited to creating parts from exiting nitinol geometries such as sheets and tubes. For these reasons, machining nitinol is a time-consuming and expensive process. Furthermore, machining complex parts out of nitinol is hardly a possibility at all.

Chapter 4

Additive Manufacturing

Additive manufacturing (AM) provides a solution to manufacturing complex nitinol components. AM technologies emerged in the mid-1980’s [165]. In earlier literature, various AM technologies may also be referred to as layer manufacturing, or solid-free form fabrication. However, in 2009, the ASTM committee F42 on Additive

Layer Technologies standardized the terminology for all layer-by-layer manufacturing as

AM.

All AM techniques utilize layer-based methods to produce 3D solid parts [166-

168]. Each layer is precisely produced according to a geometry defined in a CAD model.

Using AM enables production of 3D geometries that are not possible using traditional techniques. Features such as engineered porosity, hollow parts, curved holes and filigree structures are suddenly realizable. Furthermore, direct CAD fabrication reduces the timescale of concept-to-prototype transition.

Almost all commercial AM technologies operate according to the same general procedure. Prior to the fabrication process, layers are generated by slicing a CAD model using specialist software. Depending on the AM process, typical layer thickness can vary from 20 – 150 μm [169]. Layers are produced sequentially in the z-direction. In powder

46 based systems, a focused energy source is used to supply the required thermal energy to achieve consolidation of powder particles. The scan area is determined according to its corresponding cross-section (as calculated from the CAD model). The deposition and scanning of successive powder layers continues until the entire solid model is completed.

4.1 History of AM

Additive manufacturing began in the early1980’s with the development of rapid prototyping (RP). At first, RP was used for making non-functional prototypes from polymers [171]. It was during this time that the 3D Stereolithography (STL) file format was developed. Research conducted by Deckard [172] at the University of Austin in 1986 resulted in the development of the first SLS machines by a company called DTM in 1992.

DTM was acquired by 3-D Systems in 2001. EOS GmbH entered the SLS market in 1994 and, in 2004, they acquired the rights to all patents of DTM, UT, and 3-D Systems related to laser sintering [173]. Concurrently, Westinghouse Electric Corp. began developing

AM technology in 1988 [174]. The Westinghouse project was continued by Arcella at

Johns Hopkins University and, in 1997, Aeromet was founded. Aeromet went on to manufacture complex parts from titanium for aeropace applications using the laser engineering net shaping (LENS) process. Optomec also started commercializing a LENS system around 1997. A major breakthrough came by way of fiber laser technology in the mid-1990’s, which enabled the direct sintering of manufacturing grade metals into fully dense parts. This technology became known as selective laser melting (SLM).

A general benefit of all RP processes is the ability to manufacture complex geometries since little to no tooling is involved. This freedom of design cannot be achieved by any

47 other manufacturing technique. Furthermore, the absence of tooling reduces cost, production time and lead times during the research and development (R&D) of new and modified parts. Therefore, RP processes are well suited for industries which rely on the fabrication of parts with highly customized features, such as aerospace, automotive, biomedical and military.

4.2 Selective Laser Melting

Selective laser melting is an AM process that evolved out of SLS. Since the introduction of fiber laser technology, metallic powder particles can be fully melted to form dense parts. Some of the industrial grade materials that can be processed using SLM are 316L stainless steel, Ti-6Al-4V, Co-Cr and, most recently, nitinol.

4.2.1 Background

The first SLM system was introduced in 1999 by a German company named

Fockele and Schwarze (F&S) and utilized powder-bed technology. The system was developed in conjunction with the Fraunhofer Institute for Laser Technology and was capable of processing stainless steel [171]. In the early 2000’s, F&S entered into a partnership with MCP HEK GmbH (now MTT Technologies Group) and released the first direct metal system, the MCP Realizer 250 SLM, in 2004. Other manufactures have entered the market as competitors including Trumph, EOS, Concept Laser and Phenix

Systems.

Phenix Systems was founded in France in 2000. Phenix offers three SLM machines, the PXS, PXM, and PXL, which have small, medium, and large build capacities, respectively. The Phenix PXM (Figure 4.1) is the system used in this research and, like all Phenix products, utilizes direct sintering of parts from a metallic powder-bed using a fiber laser. Due to the fact that the powder bed is fully melted and re-solidified during the process, selective laser melting (SLM) will be the terminology used hereafter.

Figure 4.1: Phenix Systems PXM Selective Laser Melting Machine

4.2.2 Phenix Systems PXM

The PXM is a powder-based SLM machine. Prior to the fabrication process, a 3D model is drawn in CAD software. The model is then imported into Phenix Systems specialist software where the CAD model is sliced into the requisite layers. Process parameters are specified for the SLM process in Phenix Systems specialist software.

During the SLM process, a powder delivery piston feeds the powdered material vertically and a tungsten roller spreads an even layer over the fabrication piston in a horizontal and rolling motion. Average layer thickness can range from 30 μm to 150 μm depending on the material. A 300 W yttrium fiber laser (λ = 1070 nm) is used to draw the layer of sliced CAD data, melting the particles together. The laser beam is deflected by galvano mirrors, which control its move as it draws each layer. The powder particles heat up upon absorption of the laser radiation. The irradiated powder becomes fully molten and the powder particles bind to each other to produce a fully dense layer. The fabrication piston is then dropped to accommodate the next layer. The entire process is repeated

(Figure 4.2) is repeated until each layer has been manufactured. The SLM process takes place in an inert atmosphere to minimize oxidation. For some materials, nitrogen is a suitable atmosphere. However, for other materials such as titanium and nitinol, which can burn in pure nitrogen gas, an argon environment is used.

Figure 4.2: Sequence of operations of the SLM process conducted on a Phenix Systems PXM

Due to the layer-based technique, the resolution and surface quality of manufactured parts are governed by the staircase effect [170]. The relationship between the layer thickness and the orientation of a surface influence the surface quality of the part is illustrated in Figure 4.3. Thinner layers reduce the step size which results in a higher part resolution. However, fabrication by thinner layers results in longer overall processing time and is somewhat limited by the powder particle size.

Figure 4.3: Visual representation of the Staircase Effect and the principle of support structures.

Also shown in Figure 4.3 are support structures between the part and the substrate. Typically, AM parts are fabricated on scaffold or filigree structures, as opposed to directly on the substrate, to facilitate part removal. A balance must be struck when designing support structures. Due to the development of thermal stresses in the parts, the support structures must be strong enough to prevent thermal deformation, yet weak enough that they can be removed without extensive post-processing.

4.3 Existing Research on AM of Nitinol

Several groups have experimented with AM of nitinol on different machines.

However, no previous research has been done specifically addressing AM of nitinol on a machine manufactured by Phenix Systems.

Several groups have studied AM of nitinol by laser processing techniques. Work done at Ruhr-Universität Bochum in Germany [178-182], University of Liverpool in the

UK [183], and University of Applied Sciences Northwestern Switzerland [184] has all focused on SLM of nitinol on the Realizer 100. Meanwhile, research performed Purdue

University [185, 186] and Washington State University [187] has focused on AM of nitinol using the LENS technology.

Non-laser based methods which have been studied to produce porous nitinol structures include self-propagating high-temperature synthesis (SHS), hot isostatic pressing (HIP), conventional sintering, and space-holder sintering [189-192]. However, with each of these methods, pore morphology cannot be designed. In fact, pore size, distribution, and porosity can be quite difficult to control.

4.4 Process Parameters

4.4.1 Laser Power

The PXM is equipped with a 300 W ytterbium fiber laser manufactured by JK

Lasers. The laser has a beam quality of M2 < 1.2 and a Gaussian beam profile [TEM00].

The laser power on the Phenix PXM is fully adjustable up to maximum output of 300 W.

4.4.2 Beam Radius

To determine the radius of the laser beam, scan tracks were lased onto a stainless steel plate without laying down a powder layer. The width of the scan tracks were measured under an optical microscope. The radius of the laser beam was determined to be 40 μm.

4.4.3 Hatch Spacing

The overlap between neighboring scan lines is a function of the width of the scan tracks and the distance between scan tracks. The width of the scan tracks is dependent on, but not equal to, the laser beam diameter. Due to heat conduction, the width of a scan track will be wider than the laser beam diameter. Furthermore, the width a scan track will vary based on laser parameters including power and scan velocity. In general, higher laser power and slower scan velocity will increase the scan track width. The center-to- center distance between scan tracks is called the hatch spacing. If the hatch spacing is smaller than the scan track width, then there will be some percentage of overlap between the neighboring scan lines. Some overlap is required to guarantee a fully dense part.

4.4.4 Layer Thickness

By adjusting the vertical position of the build piston, the thickness of each powder layer is fully adjustable. Thinner powder layers distribute the laser energy more evenly in the z-direction. However, thicker powder layers spread more evenly and reduce the fabrication time. In the present study, 30 μm layers are used.

4.4.5 Scan Strategy

In the most basic scanning strategy, the laser trajectories are in the same direction

(Figure 4.4) and are lased in the same order, for example top-to-bottom.

Figure 4.4: Basic scan strategy showing laser trajectories all in the same direction

A slightly more complex, alternating x/y scan strategy is shown in Figure 4.5. This type of scan strategy helps to reduce local heat concentrations.

Figure 4.5: Alternating x/y scan strategy

Layer-rotation is another variable that can be introduced into the scan strategy

(Figure 4.6). Here, an alternating x/y scan strategy is shown. In step 1, the starting point is marked in the top left corner by a circle. The laser will scan the layer top-to-bottom following the directions of the arrows. In step 2, the scan strategy is rotated counter- clockwise 90 degrees and the start point moves to the bottom left corner. The scanning now proceeds right-to-left, again following the alternating directions of the arrows. The scan strategy continues to rotate 90 degrees per layer, repeating itself every four layers until all of the layers are fabricated. This type of strategy is used to further reduce local heat concentrations and to create a more homogenous solid part.

Figure 4.6: Alternating x/y scan strategy with 90 degree rotation per layer

Chapter 5

Parameter Setup, Part 1: Powder and Single Tracks

This chapter describes the powder selection and the influence of energy input into the powder bed during the SLM process. Due to the unique shape memory properties of nitinol and its strong dependence on relative Ni-Ti composition, powder selection is a careful process. Impurity contents, phase transformation, and particle size are studied.

Next, a parametric analysis begins to identify the optimal parameters for manufacturing nitinol components by SLM. The first step in identifying optimal energy input for ideal melting and re-solidification is the manufacture and analysis of single layer, single tracks.

Full parts will be analyzed in the next chapter.

5.1 Powder Selection

5.1.1 Composition

Material selection must be a very careful process due to a number of important considerations. First, a suitable composition must be chosen. In the present work, a nitinol alloy with shape memory properties was desired. Therefore, a composition

57 consisting of 50.09% nickel and 49.91% titanium (atomic %) was chosen. The nitinol ingots were purchased from Nitinol Devices & Components, Inc (Fremont, Ca).

5.1.2 Atomization

The nitinol ingots were sent directly to TLS Technik GmbH (Bitterfield,

Germany) for subsequent atomization. Due to the high temperatures required to atomize the ingot material, there is a high risk for impurity pickup during processing. Significant impurity pickup can dramatically affect the material properties and biocompatibility, both of which are extremely important to control in this study. TLS Technik GmbH was chosen because they use an atomization method specifically designed to address these issues in reactive metal powders. Their method, called Electrode Induction-melting Gas

Atomization (EIGA), combines gas atomization with a crucible free design. The removal of a crucible from the atomization process is important to eliminate contamination of the atomized material with contents of the crucible itself. During the EIGA process, shown in

Figure 5.1, a rotating electrode (the ingot material) is melted by an induction coil without any contact. A thin film of liquid metal from the melting electrode flows down a cone which forms on the exposed end of the electrode itself. The molten material is then led directly through a gas nozzle system, which atomizes the metal. The particles then solidify in an atomization tower and are collected in a powder container at the bottom of the system. This process allows for creation of very pure and spherical shaped powders, free of ceramic impurities (associated with crucible based systems), and with extremely low oxygen contents. To minimize oxygen contents, the entire process is conducted under an inert atmosphere of argon gas.

Figure 5.1: Electrode Induction-melting Gas Atomization (EIGA) [162]

5.1.3 Characterization

To ensure the quality of the powder, geometric, thermal, and chemical properties were analyzed. Figure 5.2 shows an SEM micrograph (Hitachi S-4800) of the nitinol powder which was atomized by the EIGA method by TLS Technik GmbH. The powder was separated into three particle fractions post atomization: less than 25 µm, 25 to 75

µm, and greater than 75 µm. In this study, the 25 – 75 µm particle fraction is used. This particle size is considered a good compromise of flowability, powder bed density, and impurity content. In general, finer particles possess lower flowability characteristics due to higher Van der Waals forces and higher impurity content, both as a result of a higher surface area-to-volume ratio. As shown in Figure 5.2, the particles are within the fraction

59 stated by the manufacturer and possess a spherical shape, which also improves flowability of the powder.

Figure 5.2: SEM micrograph of the nitinol powder (particle fraction 25 – 75 µm)

It is important to maintain the functional properties of nitinol after additive manufacturing and this starts with addressing the phase transformation behavior of the powder. The functional properties (shape memory behavior) are governed by a thermomechanical solid-solid phase transformation. The temperatures at which the phase

60 transformations occur is extremely dependent on material composition. The composition used in this study, Ni50.09Ti, was carefully chosen to have shape memory (as opposed to superelastic) characteristics. To evaluate the transformation temperatures, a thermal analyses was conducted according to ASTM F2004-05 by differential scanning calorimetry (DSC; PerkinElmer Diamond DSC). According to the standard, the heating and cooling rates were set at 10 K/min. The results are shown graphically in Figure 5.3 and the transformation temperatures are listed in Table 5.1. The martensite start (Ms) and martensite finish (Mf) temperatures are found to be 59 C and 26 C, respectively. The austenite start (As) and austenite finish (Af) temperatures are found to be 55 C and 93 C, respectively. Figure 5.3 does show pronounced shoulders on the high temperature side of both transformations. In some cases of Ni-rich material, these shoulders can indicate the presence of Ni-rich precipitates in the material. However, it is unlikely that Ni-rich precipitates would form in the material currently studied (Ni50.09Ti) because the Ni/Ti ratio is not high enough. Therefore, it is assumed that these shoulders are an artifact of inhomogeneous heat transfer in the powder caused by non-uniform contact between the powder and the test pan. Non-uniform contact points can result in non-uniform transformation of the entire material and cause the development of these shoulder-like features during the thermal tests.

0.3

0.2

0.1

-0.1

Heat Flow (mW/g) Flow Heat -0.2

-0.3 0 20 40 60 80 100 120 Temperature (C)

Figure 5.3: Transformation temperatures of the nitinol powder (Ni50.09Ti)

Table 5.1: Transformation temperatures of nitinol powder (Ni50.09Ti)

Ms 59 C Mf 26 C As 55 C Af 93 C

As previously mentioned, it is vital to minimize the impurity contents in AM nitinol. Not only can impurities dramatically affect the transformation behavior, but there are strict standards regulating the maximum allowable contents of nitrogen, oxygen, and carbon in nitinol devices which will be used as implantable medical devices (0.05% wt.;

ASTM F2063). To analyze the impurity contents, chemical analysis of the powder was subcontracted to Fort Wayne Metals (Fort Wayne, IN).

Two powder samples were sent to Fort Wayne Metals, Inc. for chemical analysis.

Both of the samples were from the same batch of ingot material supplied by Nitinol

Devices & Components, Inc. (Fremont, Ca) and atomized by the EIGA method by TLS

Technik GmbH (Bitterfield, Germany). One of the powder samples had a powder fraction of 25 – 75 μm and the other powder sample had a powder fraction of less than 25 µm.

Neither of the samples had been used in an SLM process prior to chemical analysis.

It was found that the nitinol powder had carbon and nitrogen contents well below the maximum allowable content of 0.05% (Table 5.2) for both powder fractions.

Furthermore, the oxygen content for the 25- 75 µm powder was also below the maximum allowable content at 0.0409% wt. However, the oxygen content for the < 25 µm powder above the maximum allowable content at 0.0629% wt. The reason for the significant increase in oxygen content in the < 25 µm powder is likely the increased surface-to- volume ratio that smaller particles possess.

Table 5.2: Chemical impurity contents of nitinol powder. Ingots provided by Nitinol Devices & Components, Inc (Fremont, Ca) and atomization by the EIGA method by TLS Technik GmbH (Bitterfield, Germany)

Carbon (% Nitrogen (%

wt) Oxygen (% wt) wt)

ASTM E1941 ASTM E1409 ASTM E1477

Powder fraction: 25-75 µm 0.0030 0.0409 0.0031

Powder fraction: < 25 µm 0.0031 0.0629 0.0032

The results of the chemical analysis clearly indicate that a powder fraction of 25 –

75 µm is favorable over a powder fraction of < 25 µm for minimizing unwanted impurities. The larger particles are also favorable for increased flowability of the material

63 during the powder deposition process for each layer. It is noted that larger particles will decrease the powder bed density. However, as 25 – 75 µm is still a relatively small powder fraction, it is believed that any effect this has during the SLM process will be negligible. Therefore, the 25 – 75 µm powder fraction is identified as the best option for

SLM of nitinol and will be used in this study.

5.2 Quality and Width of Single Tracks

The most basic structure which can be manufactured is a single layer of an individual laser scan. In this study, single tracks were manufactured with laser power ranging from 45 W to 300 W and scan velocity ranging from 0.1 m/s to 3.0 m/s. In general, higher laser power and lower scan velocities lead to increased energy input, which increasing the size of the melt and, as a result, the width of the individual track.

Figure 5.4 displays the track width as a function of scan velocity by laser power. Note that track width could not be measured for any set of parameters involving a scan velocity higher than 2.0 m/s because the energy was too low to consolidate the powder particles and form a track. The single tracks were manufactured with a 30 μm powder layer.

600 300 W 250 W 500 200 W 150 W

400 100 W 85 W 60 W 300 45 W

200 Track Width Width (µm) Track

100

0 0 0.5 1 1.5 2

Scan Velocity (m/s)

Figure 5.4: Measured single track width as a function of scan velocity for different laser powers

As expected, low scan velocities and high laser powers lead to wider single tracks. The widest track measured was nearly 600 μm (0.1 m/s, 300 W), while the smallest track measured was 66 μm (0.5 m/s, 60 W). Micrographs of selected single tracks are shown in Figure 5.5 (Meiji ML7000).

Single tracks with extremely high energy, such as the 300 W, 0.1 m/s sample, exhibit a phenomenon known as the keyhole effect. If the input energy is high enough, the powder material can be evaporated, creating a void (or a keyhole) in the powder bed.

This void absorbs the laser radiation much more effectively, leading to further increased

65 temperatures in the melt and more material evaporation. Due to the Gaussian beam profile, the maximum energy is focused in the center of the beam and, therefore, the keyhole effect is concentrated along the centerline of each track. In Figure 5.5, the keyhole effect can be seen as a trough which runs along the centerline of the single tracks produced with scan velocity of 0.1 m/s and laser powers of 200 W and 300 W. This creates a wavy surface across width of the track and makes it very difficult, if not impossible, to produce a fully dense part with these parameters.

The effect of too little energy input can also be observed in Figure 5.5. The single track produced with a scan velocity of 1.0 m/s and laser power of 100 W does not exhibit a continuous line. Here, the energy is too low to maintain a continuous melt in the powder bed. As a result, only small, globular structures are intermittently formed and adhered to the base plate. These formations are likely a result of Plateau-Rayleigh instabilities in the melt. Due to the interrupted natured of the melt, it would be impossible to manufacture a fully dense part with these parameters.

In this figure, the 200 W, 1.0 m/s track also shows signs of an unstable melt. The track is irregular and the width is somewhat uneven. However, none of the remaining tracks feature any glaring defects. On the whole, they exhibit a regular structure in both the longitudinal and transverse directions. This indicates a consistent melt with enough energy to fully melt the powder bed, but not so much to introduce the keyhole effect. In this sense, the analysis of single tracks is more of an exercise in elimination than it is in selection. After the single track analysis, a wide array of parameters are left which may be optimal for producing 3-dimensional structures. Therefore, the main benefit of the single track analysis is to identify a window of parameters which will continue to be 66 assessed by other methods. The single track analysis alone cannot be used to select an optimal set of parameters for the manufacture of dense structures.

Figure 5.5: Micrographs of SLM single tracks

5.3 Mathematical Prediction of Single Track Width

For each laser power, the data relating scan velocity to track width were fitted with a power regression (shown in Figure 5.4). A power regression fits the function:

푦 = 푎 ∙ 푥푏 6.1

Where y corresponds to track width and x corresponds to scan velocity.

The amplitude, a, and the exponent, b, of the fitting function for each laser power are listed in Table 5.3.

Table 5.3: Amplitude and exponent of the fitting functions by laser power

Laser Power (W) Amplitude, a Exponent, b

60 51.55 -0.424

85 83.589 -0.375

100 103.02 -0.348

150 123.3 -0.433

200 147.04 -0.486

250 175.67 -0.488

300 200.72 -0.47

The amplitude, a, was plotted as a function of the laser power (Figure 5.4). This relationship was fit with a linear regression:

푎 = 푚 ∙ 푃푙푎푠푒푟 + 푐 6.2

Where values for the coefficient, m, and constant, c, were determined to be 0.5727 and

32.743, respectively.

250

y = 0.5727x + 32.743 R² = 0.9724 200

a 150

100 Aplitude, Aplitude,

0 0 50 100 150 200 250 300 350 Power (W)

Figure 5.6: Amplitude, a, of the fitting function as a function of laser power

An average value of the exponent, b, was then calculated and found to be: -0.432.

Using the linear regression as the expression for amplitude, a, and the average exponent, b, a new power regression was created which comprises all of the laser powers into one expression:

−0.432 훥ℎ푠 = (0.5727 ∙ 푃푙푎푠푒푟 + 32.743) ∙ 푣푠 6.3

The calculated track widths based on the combined power regression are displayed graphically in Figure 5.7. The solid line represents the calculated track widths and the individual data points represent the measured track widths, given for reference.

700.000 300 W 250 W 600.000 200 W 150 W 500.000

m) 100 W μ 85 W 400.000 60 W

300.000

200.000 Track Width Width ( Track

100.000

0.000 0 0.5 1 1.5 2 2.5 3 Scan Velocity (m/s)

Figure 5.7: Calculated single track width as a function of scan velocity for different laser powers

The average difference between the measured track width and the calculated track width from the combined power regression is 0.8 microns with a standard deviation of 23.1 microns.

5.4 Summary

Nitinol ingots (Ni-50.09% at.) were purchased from Nitinol Devices &

Components, Inc (Fremont, Ca) and atomized by the EIGA method by TLS Technik

GmbH (Bitterfield, Germany). The powder was separated into two usable particle fractions: < 25 µm, and 25 – 75 µm. The 25 – 75 µm particle fraction was determined to be the better option based on its low impurity contents and high flowability.

The thermal behavior of the powder was characterized by DSC. The transformation temperatures were determined (Ms = 59 C, Mf = 26 C, As = 55 C, Af = 93

C), and are appropriate based on the relative Ni-Ti composition.

Using the selected powder fraction, a single track analysis was conducted to analyze the effect of process parameters (laser power, scanning velocity) on track width.

It was found that increasing laser power and decreasing scanning velocity lead to wider single tracks. Based on these results, a mathematical expression was developed to predict the width of a track based on two input variables: laser power and scan velocity. The predictor was determined to be accurate to 0.8 µm with a standard deviation of 23.1 µm.

A study of the quality of single tracks was used to develop a smaller window of parameters for further analyses. Very high energy inputs were shown to result in keyholing and very low energy inputs were not capable of creating a continuous melt.

Chapter 6

Parameter Setup, Part 2: Dense Parts

In this chapter, a parametric analysis is conducted to determine the optimal parameters for manufacturing dense nitinol components by SLM. To ensure complete solidification and the production of fully dense parts, the energy input must be high enough to fully melt the powder particles. However, there are drawbacks to operating with a higher energy input including poor surface texture and increased impurity pickup.

Therefore, it is of primary concern to find an energy input which is high enough to achieve a good melt without compromising other important characteristics.

6.1 Material Density

The challenge for powder bed additive manufacturing technologies is to produce a fully dense part from a loose powder bed. An advantage of the Phenix PXM system over other SLM systems is the compaction of the powder bed by a tungsten roller before each layer is lased. Nevertheless, a systematic approach to determine the optimal parameters for producing dense nitinol structures was devised.

Cuboid blocks were manufactured with dimensions 11.5 x 11.5 x 17.0 mm on top of support structures (Figure 6.1). A layer thickness of 30 µm was used to produce all of the samples for density analysis. The layer thickness was chosen based on the powder particle size (25 – 75 µm), desire to optimize resolution as much as possible (i.e. eliminate the staircase effect), and to facilitate remelting of the previous layer to achieve good wetting of the new layer [193]. Partially remelting previous layers is necessary to induce the epitaxial solidification required to manufacture fully dense parts. Furthermore, based on the average width of optimal single tracks and previous research [179], the hatch spacing was set to 120 µm. The color variation in each sample is due to impurity pickup during the SLM process. In general, a higher energy input leads to hotter temperatures in the melt pool which increases impurity pickup. These impurities, primarily oxygen, cause the SLM parts to exhibit a yellow hue on the surface. The SLM process is conducted under an argon atmosphere to minimize impurity pickup, however, the atmosphere cannot be perfect and some impurity pickup still occurs.

After SLM, each of the samples were removed from the plate and any remaining support structures on the parts were subsequently removed. The supports were designed to be removed easily with very little force so as not to disrupt the properties of the parts.

The top and bottom surface of each sample was precision grinded. This was done so that the top and bottom surfaces would be flat and parallel to each other, as well as perpendicular to the sides. After grinding, the sample dimensions were measured with a

Mitutoyo Vernier height gauge on a granite pedestal to determine the volume. The mass was measured on a precision scale, allowing for the density of each part to be calculated.

Figure 6.1: Cuboid blocks produced by SLM with various laser power and scan velocities for density analysis

The relative density of each part is expressed as a percentage of maximum theoretical density (assumed to be 6.45 g/cm3). The relative density is plotted as a function of scan velocity for several different laser powers in Figure 6.3. Quite predictably, there is a noticeable decrease in relative density for each power as the scan velocity is increased, corresponding to lower input energy densities. This is particularly evident in the 50 W and 100 W samples.

However, there is also a noticeable decrease in density at lower scan velocities for each power, as well. In these cases, the higher input energy leads to melt pool instability and poor surface geometry (wavy structures, Figure 6.2). This can affect the powder deposition in the subsequent layers and lead to the creation of cavities within the parts.

Figure 6.2: Poor surface quality caused by too high energy input

100

96 50 W 100 W 94 150 W 200 W 92 250 W 300 W

90 Relative Density (%) Relative 88

86 0 0.5 1 1.5 2 Scan Velocity (m/s) Figure 6.3: Relative density of SLM nitinol parts as a function of scan velocity for different laser powers

Figure 6.4 shows the relative density of the SLM nitinol parts as a function of energy density. Energy density, ωV, is calculated by the following equation (laser power,

P; scan velocity, vs; hatch spacing, Δhs; layer thickness, dt):

푃 7.1 휔푉 = 푣푠 ∙ 훥ℎ푠 ∙ 푑푡

Figure 6.4: Relative density of SLM nitinol parts as a function of energy density. Selected micrographs show the evolution of microstructure as the input energy is increased.

The energy density equation takes into consideration the amount of power from the laser and the geometric volume into which it’s directed. Included in Figure 6.4 are

76 micrographs from the cross-section of selected samples. The cuboid samples were sectioned with a Buehler IsoMet 1000 precision saw. The newly exposed cross-section was then mounted and mechanically polished to a mirror finish. Mounted samples are shown in Figure 6.5. The samples were analyzed on a Meiji ML7000 optical light microscope. The ML7000 metallurgical microscope is outfitted with a camera and all of the micrographs were taken using the software ScopePhoto.

Figure 6.5: SLM nitinol samples mounted and polished for microscopy

In Figure 6.4, it is shown that for input energies less than 40 J/mm3, there is an obvious drop off in subsequent part density. The lowest energy density tested (27.7

J/mm3) corresponds to the lowest relative density (88%). The micrograph for this sample

77 does indeed show many large pores throughout the cross-section of the part. A micrograph of a sample manufactured with energy density of 37 J/mm3 shows fewer and smaller pores, but still a relatively substantial amount.

Above 40 J/mm3, virtually all of the samples show high relative density. Among these, a micrograph is shown for one combination (P = 250 W, vs = 1.25 m/s, ωv = 55.5

J/mm3). This sample has a measured relative density of 98.3% and the micrograph confirms that it is free of any visible pores. Based on the criteria for choosing an optimal set of parameters, this is the one that is selected as best. Most importantly, it is shown that this combination of parameters is capable of producing fully dense nitinol parts. It is also significantly important that this combination results in a relatively low energy density, which will help to minimize impurity pickup during the process. A high magnification comparison of parts with the lowest measured density (ρrel = 88%, P = 100 W, vs = 1.0

3 m/s, ωv = 27.7 J/mm ) and the highest measured density (ρrel = 98%, P = 250 W, vs = 1.25

3 m/s, ωv = 55.5 J/mm ) is shown in Figure 6.6.

Figure 6.6: Comparison of pore formation in parts with lowest measured density and highest measured density at 40X magnification

It is also shown that very high energy density can result is a fully dense part (P =

3 100 W, vs = 0.1 m/s, ωv = 277.7 J/mm ). The measured relative density of this sample is

97.8% and the micrograph confirms that there are no visible pores inside the part.

However, this is undesirable because increasing energy inputs are directly related to increasing impurity pickup during the SLM process. Furthermore, while there is a general trend indicating that increasing energy input results in increasing relative density, it is not always true. As mentioned previously, high energy inputs can lead to surface tension driven instabilities in the melt pool which can create wavy surfaces. In turn, the layering process is interrupted and voids within the part can occur. This is evident in the 150 W and 200 W specimens. Above input energy density 100 J/mm3, samples manufactured at these laser powers start to show decreasing relative density. The presence of pores within

3 these samples was confirmed via microscopy (Figure 6.7, ωV = 185.2 J/mm ).

Figure 6.7: Optical micrograph of SLM nitinol manufactured with: P = 150 W, vs = 3 0.225 m/s, ωV = 185.2 J/mm

Finally, while the energy density is a useful metric to characterize trends, it is a relatively simple metric and many other factors are at play during the SLM process. Even though the set of parameters chosen as “ideal” have an energy density of 55.5 J/mm3, not every combination of parameters that has this energy density produce the same results.

The same energy density can be achieved by operating at 100 W power with a scanning velocity of 0.5 m/s. However, this results in a slightly lower measured density of 97.5% and microscopy confirms the presence of small, scattered pores in the part. Below in

Figure 6.8, this micrograph is compared directly with a micrograph from a part manufactured with the “ideal” parameter setup (P = 250 W, vs = 1.25 m/s).

Figure 6.8: Comparison of pore formation in two SLM nitinol parts manufactured with the same energy density: 55.5 J/mm3; (a) "ideal" parameter setup which shows no visible pores, and (b) visible pores ranging from 50 - 200 microns

Up to this point, all of the parts were manufactured with a hatch spacing of 120

µm. This hatch spacing was chosen based on the average width of ideal-looking single tracks. Some overlap between successive tracks is necessary to achieve complete melting and a fully dense part. To analyze the effect of hatch spacing, twelve samples were produced with different hatch distances. The variation in hatching started at 60 µm and each iteration is approximately 125% the size of the previous one. All of the parts were manufactured with otherwise constant process parameters (P = 250 W, vs = 1.25 m/s, ds =

30 µm). Figure 6.9 displays the relative density of these parts as a function of hatch spacing. It is shown that with these process parameters, hatch distances between 60 µm and 145 µm produce fully dense parts (approximately 98% relative density). Further increasing of the hatch spacing starts to yield lower density parts. This is entirely due to the lack of overlap between successive tracks. Spaces between tracks can be seen in the 81 micrographs. With a hatch spacing of 285 µm, the measured distance between hatches is approximately 85 µm. At 450 µm hatching, the distance between hatches increases to approximately 250 µm, and at 700 µm hatching, the distance between hatches increases to approximately 500 µm. Based on these results, which show that 120 µm hatch spacing produces a fully dense part, but is not too close to the point where relative density starts to decline, it is determined that this is the ideal hatch spacing to use with this parameter setup.

Figure 6.9: Effect of hatch spacing on relative density of SLM nitinol parts

6.2 Chemical Analysis

Chemical impurity pickup during SLM processing is a concern for mechanical, functional, and biocompatibility properties. High impurity pickup, particularly oxygen, leads to the formation of brittle intermetallic phases in the SLM parts which reduce the mechanical integrity and functional behavior of the SMA. Furthermore, there are strict standards regulating the maximum allowable contents of nitrogen, oxygen, and carbon in nitinol devices which will be used as implantable medical devices (0.05% wt.; ASTM

F2063).To analyze the impurity contents, chemical analysis of the powder was subcontracted to Fort Wayne Metals (Fort Wayne, IN).

Twelve SLM nitinol samples were manufactured with different combinations of laser power and scan velocity to capture a wide range of energy densities. For all twelve samples, carbon and nitrogen contents remained well below the maximum allowable value of 0.05% wt. The average carbon content in the SLM nitinol samples was 0.0053% wt. This represents a 77% increase over the carbon content in the nitinol powder. The average nitrogen content in the SLM nitinol samples was 0.0127% wt. This represents a

310% increase over the nitrogen content in the powder, but is still well below the maximum value allowed. On the other hand, the oxygen in 11 of the 12 SLM nitinol samples was found to be over the maximum allowable limit with an average content of

0.0596% wt. This represents a 46% increase over the oxygen content in the powder. The only sample that stayed within the allowable limit (manufactured with process parameters: P = 150 W, vs = 0.75 m/s) was just under the limit at 0.0493% wt. The parameters identified as ideal based on density (P = 250 W, vs = 1.25 m/s) had a slightly higher oxygen content of 0.0502% wt., just 0.0002% above the limit (Table 6.1).

Table 6.1: Impurity contents measured in powder and SLM nitinol part (P = 250 W, vs = 1.25 m/s)

Carbon (% Nitrogen (%

wt) Oxygen (% wt) wt)

ASTM E1941 ASTM E1409 ASTM E1477

Nitinol powder 0.0030 0.0409 0.0031 (25-75 µm fraction)

SLM nitinol part 0.0035 0.0502 0.007

(P = 250 W, vs = 1.25 m/s)

Despite an oxygen content that is slightly above the maximum allowable limit, this is a favorable result for the parameter setup: P = 250 W, vs = 1.25 m/s. Of the twelve samples tested, this oxygen content is the second lowest that was measured and is only

0.0002% above the limit. Furthermore, the oxygen content only increased 23% from that which was measured in the powder. Therefore, it is fully expected that with stricter quality control of ingot and powder material processing, the oxygen content in the SLM nitinol can be kept well within the allowable range with these process parameters.

In general, higher energy densities led to greater pickup of impurities. This trend is shown for oxygen in Figure 6.10.

0.09

0.08

0.07

0.06

0.05 Oxygen Content (% wt) wt) (% Content Oxygen 0.04 0 50 100 150 200 250 300 Energy Density (J/mm3)

Figure 6.10: Oxygen content in SLM nitinol parts manufactured at various energy densities

6.3 Summary

A parametric analysis of SLM process parameters on part density was conducted.

Cuboid blocks were produced with a wide range of parameters. Laser power was varied from 50 W to 300 W, and scan velocity was varied from 0.1 m/s to 1.875 m/s. The layer thickness and hatch spacing were kept constant at 30 µm and 120 µm, respectively. The relative density for each part was measured and pore formation was confirmed via microscopy. In general, it is shown that relative density increases with increasing energy density (lower scan velocity and/or higher laser power). However, it is also shown that high energy inputs can also lead to pore formation within the part, thus reducing its density.

The effect of hatch spacing on part density was studied and 120 µm was found to be optimal. 85

A chemical analysis was conducted to investigate the effect of impurity pickup.

All of the SLM nitinol parts were well within the limits for carbon and nitrogen content.

It was shown that the oxygen content was just over the allowable limit (+0.0002%), but only increased 23% over the content in the powder. Therefore, with stricter control of the impurities in the powder, the oxygen content can be kept well within standard as well.

Based on this study, the optimal parameters for SLM processing of nitinol on a

Phenix Systems PXM machine have been determined (Table 6.2). A single track is shown in Figure 6.11 and has a measured width of 163 µm. The model presented in

Section 5.3 predicts a track width of 159.8 µm with these parameters. Figure 6.12 shows the top surface of an SLM nitinol part produced with the optimal parameter setup. With this setup, each track overlaps 36% with the adjacent track.

Table 6.2: Optimal parameters for SLM processing of nitinol on a Phenix Systems PXM Laser Power 250 W

Scan Velocity 1.25 m/s

Hatch Spacing 120 µm

Layer Thickness 30 µm

Figure 6.11: Micrograph of SLM nitinol single track produced with optimal parameter setup

Figure 6.12: Micrograph of surface of SLM nitinol part produced with optimal parameter setup

Chapter 7

Thermal, Mechanical, and Functional Properties of SLM Nitinol

This chapter describes the thermal, mechanical, and functional behavior of SLM nitinol structures. Nitinol exhibits unique thermomechanical behaviors including a solid- solid phase transformation from a low temperature martensitic phase to a high temperature austenitic phase, as well as twinning in the martensitic state. To show that

SLM nitinol structures exhibit these properties, a series of tests were performed. First, the phase transformation was characterized by differential scanning calorimetry (DSC).

Compression testing was done in the martensitic state to observe the stress-strain relationship during detwinning and to analyze the ultimate strength. Finally, the cyclic behavior of the SLM nitinol was investigated to prove the shape memory effect.

7.1 Transformation Characteristics

The thermal behavior of SLM nitinol parts was measured on a PerkinElmer

Diamond DSC. Cylindrical SLM parts (diameter = 4.5 mm, height = 10 mm) were manufactured in the vertical direction with process parameters: P = 250 W, vs = 1.25 m/s,

Δhs = 120 µm, ds = 30 µm. A cross-sectional slice approximately 500 µm thick was removed from the samples with a precision saw (Buehler Isomet 1000) for thermal 89 analysis in the DSC. As shown in Figure 7.1, the SLM process results in slightly increased transformation temperatures when comparing the SLM nitinol parts with powder. For the SLM part, the martensite start (Ms) and martensite finish (Mf) temperatures were found to be 62 C and 35 C, respectively. The austenite start (As) and austenite finish (Af) temperatures were found to be 65 C and 92 C respectively.

0.3 SLM nitinol 0.2 Powder

0.1

-0.1

-0.2 Heat Flow (mW/g) Flow Heat

-0.3 0 20 40 60 80 100 120 Temperature (C)

Figure 7.1: Transformation temperatures of SLM nitinol part and nitinol powder

The transformation temperatures of the SLM nitinol part and powder are compared directly in Table 7.1. The first thing which is noticed is that the shoulder on the high temperature side of the DSC curves for the powder material do not occur for the

SLM part. This further substantiates the theory that these shoulders are purely a function of inhomogeneous heat transfer between the DSC pan and the powder due to non- uniform contact. Next, it is seen that there is a shift in transformation temperatures after

90 the SLM process towards higher values. At the transformation peaks, the temperatures are shifted higher by 16 C in both the heating and cooling directions. However, because the transformation region becomes skinnier for the SLM part, the temperature increase is not uniform for each transformation. For example, Ms only increases by 3C, while Mf increases by 9 C. Furthermore, As increases 10 C after the SLM process while Af actually decreases by 1 C.

Table 7.1: Transformation temperatures of SLM nitinol part and nitinol powder Powder SLM Part Ms 59 C 62 Mf 26 C 35 As 55 C 65 Af 93 C 92 Mpeak 34 C 50 Apeak 66 C 82

The increase in transformation temperatures after SLM can be attributed to nickel evaporation during processing. Nickel has a lower evaporation temperature than titanium and, therefore, a higher tendency to evaporate during the process [178-181]. As discussed in Chapter 3, the transformation temperatures of nitinol are extremely sensitive to alloy composition [147]. For this reason, nickel evaporation and the resulting shift in transformation temperatures should be considered when selecting powder for manufacturing nitinol components by SLM.

7.2 Compression Testing

The mechanical properties of SLM nitinol were investigated by compression testing on an Instron 5569 series tension-compression test system using Bluehill software.

All samples were compressed at a rate of 0.5 mm/min. Strain was measured with an

Instron strain gauge and the system was outfitted with a ± 50 kN load cell.

For fracture testing, cylindrical SLM nitinol samples (diameter = 4.5 mm, height

= 10 mm) were manufactured in the vertical direction with process parameters: P = 250

W, vs = 1.25 m/s, Δhs = 120 µm, ds = 30 µm. The samples were reduced in height to 7 mm by removing 1.5 mm from the top and bottom of the sample on a precision saw. The contact faces were then hand polished to ensure uniform contact between the specimen and the compression plates. No modifications such as grinding or turning were done to the sides of the SLM nitinol samples prior to compression testing. For comparison, compression samples were produced from conventionally manufactured nitinol ingots.

The ingot material used was taken from the same batch which was atomized to obtain the powder used for the SLM samples.

Prior to compression testing, the samples were heated to 120 C in an oven for 20 minutes, which is 30 C above Af, to reach a fully austenitic state and then cooled to 0 C in a freezer for 20 minutes to guarantee a state of twinned martensite throughout the material. Testing was then performed at 21 C, where the samples remain fully martensite, well below the austenite start temperature of 65 C.

4000 SLM nitinol 3500 Conventional nitinol

3000

2500

2000

1500 Stress (MPa) Stress

1000

500

0 0 5 10 15 20 25 30 35 Strain (%)

Figure 7.2: Stress-strain curve for SLM and conventional nitinol in the martensitic state

Figure 7.2 shows a stress-strain curve for SLM nitinol and conventionally produced nitinol in the martensitic state. In both samples, the material is initially deformed according to its martensite modulus, measured to be approximately 31 GPa. At approximately 150 MPa, the material exhibits softening at the onset ofA martensitic detwinning. An inclined stress plateau continues until the material is fully detwinned and starts to compress again according to its martensite modulus. The SLM nitinol appears to complete martensitic detwinning earlier than the conventional sample. Furthermore, it seems that plastic deformation occurs at lower stress values in the SLM sample compared to the conventional sample. The variances could be due to different microstructures in the

SLM and conventional nitinol. Higher residual stresses in the SLM nitinol which affect its ductility could also have an effect. A definitive conclusion cannot be made without a comprehensive study of the microstructural behavior during loading of each specimen.

However, since nitinol is a functional material and is not subjected to these high strains in practice, this is not of primary concern in this study. The plateau stresses and strains are considerably more important in characterizing the quality of the nitinol part. The absence of a flat detwinning plateau in compression is well known and reported for conventional

[199] and SLM nitinol [200]. The reason for this is a high density of dislocations during compression, which quickly induces hardening of the material. Fracture occurred in the

SLM nitinol at about 3000 MPa and 29.5% strain. For the conventional sample, fracture occurred at approximately 3400 MPa at 31.5% strain.

In every sample, fracture occurred at an angle of 45º to the loading direction, corresponding to the direction of highest shear stress (Figure 7.3). Haberland et al. reported fracture in SLM nitnol samples around 3300 MPa at approximately 40% strain for SLM nitinol samples and 3900 MPa at 47% strain for conventional nitinol [195].

Figure 7.3: Fracture of SLM nitinol samples at 45º with respect to the loading direction

7.3 Cyclic Testing

To evaluate the functional properties of SLM nitinol, cylindrical test samples

(diameter = 6.25 mm, height = 12 mm) were manufactured in the vertical direction with process parameters: P = 250 W, vs = 1.25 m/s, Δhs = 120 µm, ds = 30 µm. Using the same method as described in the previous section, the parts were reduced in height to 10 mm with a precision saw and the contact faces were hand polished to ensure even contact.

The sides of the test samples were not subjected to any form of post processing before mechanical testing. Cyclic tests were performed up to two stress levels: 300 MPa and 900

MPa. These stress levels were chosen based on the compression data presented in Figure

7.2. The 300 MPa stress level corresponds to the end of the plateau region (complete detwinning), and the 900 MPa stress level corresponds to a point shortly before the end of the elastic region of the detwinned martensite. 95

Before each cycle, the length, Ln, and diameter, Dn, were measured. The samples were then heated to 120 C, corresponding to 30 C above Af. This heating cycle allows for the material to fully transform into austenite and results in the shape recovery of reversible deformation. The samples were then cooled to 0 C, which is 30 C below Mf, to revert the material back to its twinned martensite state. Each measurement of length, Ln, was used to determine the strain in the nth cycle and the irreversible strain from cycle n-1.

The tests were conducted at 21 C, well below the austenite start temperature of 65 C.

The cyclic mechanical behavior of SLM nitinol, evaluated to 300 MPa, is shown in Figure 7.4. The cyclic behavior evaluated to 900 MPa is shown in Figure 7.5. Shape recovery is found to be very similar to that of conventionally produced nitinol, as well as to previously reported studies on SLM nitinol, for both cases. It is shown that there is a degradation of the shape memory effect, particularly in first few cycles. This is widely reported for all nitinol shape memory alloys [195], including those conventionally manufactured.

Figure 7.4: Stress-strain curves of shape memory cycling to 300 MPa

The cumulative irreversible strain during cyclic loading is given for both loading conditions in Figure 7.6. Most of the irreversible strain is accumulated during the first few cycles. Then, as the cycles continue to progress, the shape memory very clearly begins to stabilize and less irreversible strain is accumulated. For the SLM nitinol specimen that was loaded to 300 MPa (Figure 7.4), 0.6% irreversible strain occurred during the first cycle. By the eighth cycle, a total of 2.1% irreversible strain had accrued.

However, from the ninth to the fifteenth cycles only 0.31% more irreversible strain occurred, resulting in a total irreversible strain of 2.42% for 15 cycles.

Figure 7.5: Stress-strain curves of shape memory cycling to 900 MPa

As expected, cyclic loading to 900 MPa (Figure 7.5) resulted in much higher irreversible strains. However, as shown in Figure 7.6, most of the irreversible strain occurs during the first few cycles and then the shape memory behavior stabilizes. In the first cycle alone, there was 3.6% irreversible strain. By the eighth cycle, a total of 5.64% irreversible strain had accrued, after which the shape memory behavior became a lot more regular. From the ninth to the fifteenth cycle only 0.53% more irreversible strain occurred, resulting in a total irreversible strain of 6.17% for fifteen cycles. Haberland et

98 al. reported similar trends for the behavior of SLM and conventional nitinol in cyclic testing [195].

7 σ-max = 300 MPa

6 σ-max = 900 MPa

1 Cumulative Irrev. Strain (%) Strain Irrev. Cumulative

0 0 2 4 6 8 10 12 14 16 Cycle Number, n

Figure 7.6: Cumulative irreversible strain in SLM nitinol during cyclic testing to different stress levels

7.4 Summary

SLM nitinol structures were manufactured with cylindrical geometries for thermal, mechanical, and functional analysis. The optimal parameter setup determined in

Chapter 6 was used to manufacture all of the test samples. For thermal analysis, a thin cross-sectional slice was removed from the middle of a samples. For compression testing, the top and bottom surfaces of each sample were polished to ensure uniform loading with the compression plates. No post-processing was done to the sides of the test samples before compression.

Thermal analysis was conducted on a PerkinElmer Diamond DSC. It was found that the transformation temperatures are in the shape memory regime as expected based on the Ni-Ti composition. Furthermore, the transformation temperatures are generally shifted to higher values, likely due to nickel evaporation in the material during laser processing.

The stress-strain behavior of SLM nitinol in compression is the very similar to conventionally processed nitinol. Fracture occurred in the SLM nitinol at about 3000

MPa and 29.5% strain. In the conventional nitinol sample, fractured occurred at 3400

MPa and 31.5% strain. In each sample, fracture occurred at an angle of 45º to the loading direction, corresponding to the direction of highest shear stress.

The functional behavior of SLDM nitinol was studied by cyclic testing to two stress levels: 300 MPa and 900 MPa. Considerable irreversible strain was accumulated during the first few cycles. However, it is was shown that after a number of cycles, the shape memory behavior stabilizes and very little irreversible strain is incurred.

100

Chapter 8

Designed Porosity in Nitinol Structures

In this chapter, nitinol structures with engineered porosity are studied. Engineered porosity in nitinol parts can be used to modulate the effective stiffness of the structure.

For use in implants, this can be used to reduce the stiffness to a range which is similar to bone. Furthermore, by designing the pores by CAD, the user maintains exact control over the size, shape, and distribution of pores in the structure. SLM nitinol parts were designed with the properties of bone in mind. These parts were mechanically tested in compression and compared to current implant metals.

8.1 Review of Metallic Implant Requirements

As detailed in Chapter 2, the metals commonly used in long-term implants are much stiffer than bone. When the bone and implant are loaded, the stiffness mismatch results in an effect called stress shielding. Stress shielding refers to the redistribution of load from the bone to the implant. Due to its high stiffness, the implant absorbs the stress without undergoing much strain. This also reduces the strain in the bone which inhibits bone remodeling (the biological process which maintains the structural and mechanical integrity of bone). In the absence of applied strain, the bone surrounding the implant 101 naturally degrades, which can ultimately lead to implant failure. Due to its lower stiffness, nitinol is investigated as a more biomechanically friendly alloy than the traditional materials, stainless steel and titanium.

8.2 Computer-aided Design of Porosity

Porous structures were designed with biomorphic geometries to meet two major goals: reduce the apparent stiffness of the structure to a value similar to bone and promote bone-ingrowth. To promote bone-ingrowth, the scaffold architecture is designed with interconnecting pores. A unit cell, shown in Figure 8.1 (a), was created. This is a simple structure which can be repeated in 3-dimensions to obtain a porous part. The strut diameter and length can be modulated to change the porosity of the part.

In this research, three porous geometries based on the unit cell were studied. Each of the geometries featured 128 unit cells, repeated in a 4x4x8 pattern. The dimensions of each unit cell and the resulting pore sizes are given in Table 8.1.

102

Figure 8.1: (a) Unit cell of porous design with strut diameter, d, and length, l; and (b) representative structure made by repeating the unit cell in the X, Y, and Z directions [197]

Table 8.1: Dimensions of unit cells for porous structures and resulting pore size [197] Porosity by volume Strut diameter (mm) Strut length (mm) Pore Size (mm) 32% 1.4 1.0 0.6 45% 1.2 1.0 0.8 58% 1.0 1.0 1.0

103

8.3 Compression Testing

All of the porous samples were manufactured on the Phenix Systems PXM with the ideal parameter setup (P = 250 W, vs = 1.25 m/s, Δhs = 120 µm, ds = 30 µm). The porous SLM nitinol structures are shown in Figure 8.2 and micrographs of the pore structures are shown in Figure 8.3.. From left to right, the structures’ porosity by volume is: 58%, 45%, and 32%.

Figure 8.2: Porous SLM nitinol structures; from left to right: 58% porosity, 45% porosity, 32% porosity

104

Figure 8.3: Micrographs of porous SLM nitinol structures; from left to right: 58% porosity, 45% porosity, 32% porosity

To evaluate the mechanical properties of the porous SLM nitinol structures, the samples were tested in compression on an Instron 5569 series tension-compression test system using Bluehill software. All samples were compressed at a rate of 0.5 mm/min.

Strain was measured with an Instron strain gauge and the system was outfitted with a ±

50 kN load cell.

The results of the compression tests are shown in Figure 8.4 and Figure 8.5. In

Figure 8.4, the stress-strain curves are plotted all the way to 12% strain. Included in the plot are stress-strain curves for dense SLM nitinol, SLM nitinol with three different levels of porosity (32%, 45%, and 58%), and conventional metal implant materials, 316 stainless steel and titanium (Ti-6Al-4V). Young’s modulus of 316 stainless steel and titanium are assumed to be 193 GPa and 116 GPa, respectively [198]. They are only plotted out to their respective yield strength, given as 190 MPa for stainless still and 795

MPa for titanium [198].

105

1800 SLM NiTi SLM NiTi: 32% Porosity 1600 SLM NiTi: 45% Porosity SLM NiTi: 58% Porosity 316 Stainless Steel 1400 Titanium (Ti-6Al-4V)

1200

1000

800

Stress (MPa) Stress 600

400

200

0 0 2 4 6 8 10 12 Strain (%)

Figure 8.4: Stress-strain curves to 12% strain for SLM nitinol with various levels of porosity compared to conventional implant materials 316L and Ti-6Al-4V

Bone remodeling occurs as a strain response in bone [22-28]. Strains ranging from 500 – 3000 microstrain (.05% - 0.3%) are required to stimulate bone remodeling.

Furthermore, the ultimate strength of bone is typically between 120 – 200 MPa [199-

201]. The stresses in bone cannot be measured directly, but it is known that the stress- strain relationship is not linear [201]. Bone is stiffer at smaller loads and strains than at larger ones. Some studies have been conducted to estimate the elastic modulus of bone in compression and estimate it in the range of 11 – 18 GPa [199]. Based on the estimated elastic modulus and the remodeling thresholds, it is estimated that normal stresses during daily activity are in the range of 8 – 30 MPa. 106

As shown in Figure 8.4, and closer in Figure 8.5, stainless steel and titanium are so stiff that these levels are strain would be unattainable at such low stress levels. Nitinol is a desirable alternative due its low stiffness and high strength. Figure 8.4 shows all of the porous SLM nitinol configurations have an ultimate strength well above that of bone

(~200 MPa).

200 SLM NiTi 180 SLM NiTi: 32% Porosity SLM NiTi: 45% Porosity 160 SLM NiTi: 58% Porosity 140 316 Stainless Steel Titanium (Ti-6Al-4V) 120

100

60 Stress (MPa) Stress 40

0 0 0.5 1 1.5 2 2.5

Strain (%) Figure 8.5: Stress-strain curves to 2.5% strain for SLM nitinol with various levels of porosity compared to conventional implant materials 316L and Ti-6Al-4V

In Figure 8.5, the strain and stress values are reduced to the maximum ultimate strain and stress values for bone as reported in literature: 2.5% and 200 Mpa, respectively

[22-28, 199-201]. Here, it is shown that the three SLM porous nitinol samples operate

107 well within this range. Furthermore, porous SLM nitinol samples exhibit a stiffer linear region at smaller strains with elastic moduli that match the range reported for bone (Table

8.2). The linear region approximately captures the region in which adaptive bone remodeling is stimulated (.05% - 0.3% strain). Following this region, the porous SLM nitinol samples exhibit softening due to martensitic detwinning. This matches the behavior of bone which also softens at higher strains [201].

Table 8.2: Elastic Modulus for typical implant materials, SLM nitinol at various porosities and bone

Material Elastic Modulus (GPa) 316 Stainless Steel 193 Titanium (Ti-6Al-4V) 116 SLM nitinol 38.5 SLM nitinol: 32% Porosity 16.3 SLM nitinol: 45% Porosity 13.7 SLM nitinol: 58% Porosity 9.9 Bone [199] 11-18

The most porous SLM nitinol sample (58% porosity) was tested until fracture

(Figure 8.6). The ultimate strength was found to be approximately 580 MPa at 14% strain, well above the maximum in bone. Furthermore, the fracture was found to occur at a 45 degree angle with respect to the loading direction, corresponding to the direction of highest shear stress.

108

Figure 8.6: (a) Stress-strain curve of SLM nitinol with 58% porosity up to fracture and (b) fracture at 45º with respect to the loading direction

8.4 Summary

To improve the outcome of long-term metallic implant use, the mechanical properties of implants need to better match those of bone. Studies have estimated that the elastic modulus of bone in compression is in the range of 11 – 18 GPa, strains ranging from 500 – 3000 microstrain (.05% - 0.3%) are required to stimulate bone remodeling, and the ultimate strength of bone is in the range of 120 – 200 MPa [199-202]. In young adults, bone fractures at approximately 25,000 microstrain (2.5%) [28].

Nitinol structures with engineered porosity were designed and manufactured by

SLM to meet these requirements. It is shown that pore size, shape, and distribution can be customized by CAD and manufactured by SLM. In this study, porous nitinol structures were manufactured with three different porosities (32%, 45%, and 58%) and featured interconnecting pores ranging from 0.6 mm to 1.0 mm in size.

109

The porous nitinol structures were mechanically tested in compression and each showed high strength and low stiffness characteristics. The ultimate strength of each porous structure far exceeds that of bone, with the most porous structure exhibiting an ultimate strength of approximately 580 MPa. It was shown that the stiffness of SLM nitinol structures can be modulated based on the level of porosity to match the stiffness of bone. The elastic moduli of porous nitinol samples produced by SLM ranged from 9.9 –

16.3 MPa. By altering the pore size, the elastic modulus can be adjusted to match any number.

110

Chapter 9

Conclusions and Recommendations for Future Work

This research was aimed at developing a process to manufacture porous and stiffness-tailored nitinol components for use in long-term metallic implants. A comprehensive literature review revealed that there is a large unmet need for new materials and manufacturing techniques to improve existing implant technologies.

Current implant technologies are expected to restore proper biomechanical function in the body, yet utilize materials that possess drastically different mechanical properties than the biological materials they replace. Furthermore, these implants are introduced into highly unique environments, yet are mass designed with little regard for the individual patients’ needs.

To meet this unmet need, nitinol is identified as a more suitable material for use in long-term metallic implants than those which are currently used. It features high strength, low stiffness, high recoverable strain, and good biocompatibility. To manufacture the nitinol implants, an additive process called selective laser melting (SLM) is utilized. SLM circumvents current issues with manufacturing nitinol components (i.e. conventional machining of nitinol is incredibly difficult) and opens up possibilities for patient-specific

111 implant design. To this end, the manufacturing of nitinol components by SLM on a

Phenix Systems PXM was studied.

9.1 Specific Conclusions

 A single track analysis was conducted to analyze the effect of process parameters

(laser power, scanning velocity) on track width. It was found that increasing laser

power and decreasing scanning velocity lead to wider single tracks.

 High energy inputs were shown to result in uneven surfaces across the width of

single tracks due to the keyhole effect. Low energy inputs were not capable of

sustaining a continuous melt.

 Single track width can be accurately predicted based on two input variables: laser

power and scan velocity. The predictor was determined to be accurate to 0.8 µm

with a standard deviation of 23.1 µm.

 In general, it is shown that relative density of 3-dimensional structures increases

with increasing energy density (lower scan velocity and/or higher laser power).

However, it is also shown that too high energy inputs can also lead to pore

formation within the part, thus reducing its density.

 The optimal parameters for SLM processing of nitinol on a Phenix Systems PXM

are determined to be: Laser power = 250 W, scan velocity = 1.25 m/s, hatch

spacing = 120 µm, and layer thickness = 30 µm.

 Using the optimal parameter setup, relatively density of SLM nitinol parts was

measured to be 98%. Internal porosity was examined by a cross-sectioning

technique, and no visible pores were discovered. 112

 Single tracks produced with the optimal parameter setup have a measured track

width of 163 µm. The model presented in Section 5.3 predicts a track width of

159.8 µm with these parameters.

 The transformation temperatures in SLM nitinol are generally shifted to higher

values compared to the powder. This is likely due to nickel evaporation in the

material during laser processing.

 The stress-strain behavior of SLM nitinol in compression is very similar

compared to conventional nitinol. Fracture occurred in the SLM nitinol at about

3000 MPa and 29.5% strain. Fracture in the conventional nitinol occurred at 3400

MPa and 31.5% strain. In each sample, fracture occurred at an angle of 45º to the

loading direction, corresponding to the direction of highest shear stress.

 The functional behavior (shape memory effect) of SLM nitinol was studied by

cyclic testing to two stress levels: 300 MPa and 900 MPa. Considerable

irreversible strain was accumulated during the first few cycles. After fifteen

cycles, the shape memory behavior is stabilized and very little irreversible strain

is incurred.

 Porous nitinol structures were manufactured with three different porosities (32%,

45%, and 58%) and featured interconnecting pores ranging from 0.6 mm to 1.0

mm in size. This demonstrates the ability to customize pore size, shape, and

distribution by CAD.

 The elastic moduli of porous nitinol samples produced by SLM ranged from 9.9 –

16.3 MPa, in the range of the elastic modulus of bone.

113

 The ultimate strength of each porous structure far exceeds that of bone, with the

most porous structure exhibiting an ultimate strength of approximately 580 MPa.

9.2 Recommendations for Future Work

In this research, nitinol was processed on a Phenix Systems SLM machine for the first time. Process parameters were studied and optimized for the production of fully dense nitinol parts. Significant results were obtained and are previously detailed. Based on these, recommendations for future work within the overall scope of this research are made.

In this study, a nitinol composition was used which results in shape memory behavior (Ni50.09Ti). This was chosen because the shape memory effect is not dependent on any post-processing procedures. Therefore, the functional properties of SLM nitinol parts (i.e. shape memory behavior) could be investigated directly after SLM processing.

This allows for a direct comparison between process parameters and functional outcome.

On the other hand, significant post-processing is required to establish superelastic behavior, even in conventionally processed nitinol [132]. These post-processing techniques involve aging treatments which are used to precipitate Ni-rich phases of type

Ni4Ti3. The presence of Ni4Ti3 phases results in precipitation hardening of the material and impedes dislocation movement. Both effects serve to increase the yield stress in the material so that irreversible deformation is reduced. Ni4Ti3 particles also provide nucleation sites to support the martensitic transformation [155, 203-204].

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Therefore, it is recommended that the next phase of this research be focused on the development of SLM superelastic nitinol components. The optimal process parameters identified in this research can be directly translated to manufacture nitinol components with superelastic compositions. Therefore, future work needs only to address post-processing routes to establish the necessary microstructure for superelatic behavior.

The development of engineered porosity in SLM nitinol parts was also studied in this research. Until now, no other research had reported the modulation of stiffness in nitinol components by CAD-designed porosity. In this research, all of the nitinol components featured regular porosity. It is recommended that future studies analyze the effect of irregular CAD-designed porosities including asymmetrical pore sizes, shapes, and distributions.

Furthermore, based on the direct-from-CAD manufacturing technique, this research demonstrates the possibility for patient-specific implant design. It is recommended that future research focus on the development of patient-specific implants from medical image data that captures the patients’ bone geometry and stiffness.

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Chapter 10

Business Plan

10.1 Executive Summary

Ortho3d has developed a novel, biomechanically accurate total knee replacement

(TKR). The Ortho3d TKR utilizes a smart material, nitinol, as a better alternative to traditional metallic alloys. Furthermore, the Ortho3d TKR is produced through an additive manufacturing, or 3D printing, method called selective laser melting (SLM). The use of SLM to manufacture the Ortho3D TKR means the implant is highly customizable on a patient-specific basis. Knee replacement surgery is performed in patients that have severe knee pain and disability from rheumatoid arthritis, osteoarthritis, or traumatic injury. The Ortho3d TKR aims to fill a gap in the current market by:

 Utilizing a material and a manufacturing process for our implants that represent a

radical innovation in the medical device industry.

 Proving a need to the ageing, baby-boomer population which provides a

sustainable domestic market.

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Total knee replacements (TKR) are one of the most popular implant procedures performed in the United States. TKRs are used to relieve the knee joint of pain and disability, typically caused by osteoarthritis. Other major causes include meniscus tears, cartilage defects, and ligament tears.

Last year, more than 600,000 knee replacements were performed in the United

States. Additionally, over 50,000 revision surgeries were performed. The primary cause for revision surgeries is implant failure due to aseptic loosening femoral component.

Current metal materials used in the femoral component are much stiffer than bone. This leads to stress-shielding effects which causes osteoporosis of the surrounding bone and, eventually, aseptic loosening and implant failure.

The demand for primary total knee replacements and revisions are both projected to increase by over 600% by 2030 resulting in 3.48 million TKRs and over 250,000 revisions. With more young people undergoing total knee athroplasties and average life expectancy rising, it’s possible that the number of revision surgeries will increase even more. To increase the success rate of TKRs and reduce revision surgeries, it is necessary to develop better biomaterials for use in the femoral component.

The Ortho3d TKR is estimated to gain about 1% of the US market share of the total knee replacement market in the 7th year of existence and then acquire 0.25% each subsequent year. The cost of the Ortho3d TKR.in the United States will be $9000. With regards to the cost of device production, more than 600,000 demands per year for the total knees, and the market share of 1% in the 7th year, 1.25% in the 8th year, and 1.5% in the 9th year, we are projecting an annual sale of $54 million, $67.5 million, and $81 million,

117 respectively. Based on initial funding estimates, we expect that in the Q2 of our sixth year, our company will reach a financial break-even point.

10.2 Opportunity Rationale

There are two primary reasons why Ortho3d has excellent business potential.

First, the material and the manufacturing technique utilized in our implants represent a radical innovation in the medical device industry.

Total knee implants utilize multiple materials in their construction. Typically, the tibial insert and patellar component are made of plastics such as ultra-high molecular weight polyethylene (WHMWPE) or cross-linked polyethylene (XLPE), while the femoral component and the tibial tray are made of metals such as titanium alloys, stainless steel or cobalt-chromium-molybdenum. The opportunity for innovation is in the redesign metallic components.

The main failure mode of knee implants, and cause for revision surgeries, is aseptic (non-infected) loosening of one of the metallic components. The main causes of aseptic loosening are excessive wear between articular surfaces, stress shielding of the bone by the prosthesis, and development of a soft tissue at the bone-implant interface.

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The femoral component interfaces with bone from the upper side and articulates against the polyethylene insert from the lower side. Therefore, aseptic loosening of the femoral component is involved in all three causes of failure.

Nitinol possesses an advantageous combination of properties which can improve the femoral component of TKRs including high strength, high corrosion and wear resistance, good biocompatibility and osseointegration and similar elastic modulus to bone. However, fabrication of nitinol components, especially porous, is incredibly difficult using traditional machining techniques. Therefore, a technological barrier has existed which has kept nitinol from being utilized.

With the development of the technology in this dissertation, that barrier can be overcome. The use of selective laser melting (SLM) techniques will allow for the rapid manufacturing of porous nitinol components which can be used in total knee implants. In addition to realizing nitinol as an implant material, SLM can make implants more patient- specific leading to even better clinical outcomes. SLM begins with a CAD model of the component to be created, which can be derived directly from CT scans of the patient.

None of these techniques are currently being utilized in the implant industry.

Second, there is an ageing population which provides a sustainable domestic market. In 2005, the worldwide total orthopedic market (not just joint replacement) had a market size of $25.6 billion. Of this, knees represented 18% of the market with a 14% growth rate. As of 2005, the US represented 63% ($16.1 billion) of the worldwide orthopedic market, making it the largest market in the world. Europe and Japan are the next two largest orthopedic markets [204].

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In 2010, approximately 600,000 TKR procedures were performed in the United

States at a cost of approximately $15,000 per procedure ($9 billion aggregate) [204-211].

Additionally, over 60,000 revision surgeries were performed in 2010, an additional surgery to fix problems. These figures are projected to grow dramatically over the next

20 years. The number of TKRs in the United States are projected to increase by 673%, resulting in 3.48 million annual procedures, by 2030. Revision surgeries are projected to double by 2015 and increase by 601% by 2030. Approximately 61% of surgeries are performed on patients older than 65.

The ageing population alone provides a huge, sustainable domestic market. The baby boomer generation, comprising 77 million people, is just entering senior age.

Between 1991 and 2010, 3.27 million patients ages 65 and older had total knee replacements. During that same time, 318,563 of the same demographic had knee revisions. Among Medicare patients, the number of annual TKRs increased 162% from

1991 to 2010, from 92,260 to 248,802. The number of revisions also increased a similar amount in Medicare patients from 9,650 to 19,871. Total knee replacements are not only increasing as a function of the growing population. The number of these surgeries being performed has nearly doubled per capita for Medicare enrollees, from 31.2 procedures per 10,000 in 1991 to 62.1 procedures per 10,000 in 2010 [212].

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10.3 The Company

After studying the market for total knee arthroplasties, we have identified a gap in the marketplace that can be partially filled with our product, the Ortho3d TKR. Our competitors use one of three metals materials in their designs for total knee replacements.

They either use 316 stainless steel, titanium (Ti-6Al-4V) or cobalt chrome. Each of these materials can be 5 – 15 times stiffer than the surrounded bone tissue. During long-term fixation, this severely disrupts the physiomechanical processes which regulate bone remodeling. As a result, surrounding bone tissue degrades which leads to implant loosening and, ultimately, catastrophic failure. The Ortho3d TKR offers a better solution by using the shape memory alloy, nitinol. Nitinol is identified as a more suitable material for use in long-term metallic implants due to its features of high strength, low stiffness, high recoverable strain, and good biocompatibility.

Additionally, an additive manufacturing processing called selective laser melting

(SLM) is identified as a better processing route than traditional methods. SLM circumvents current issues with manufacturing nitinol components (i.e. conventional machining of nitinol is incredibly difficult) and opens up possibilities for patient-specific implant design. The radical innovations of utilizing nitinol and additive manufacturing will begin a new era of total knee replacement.

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10.4 The Product

The Ortho3d TKR is a total knee replacement for patients whose knee has been severely damaged by arthritis or injury. The Ortho3d TKR is for patients whose knee pain cannot be eliminated by non-surgical treatment methods such as medications or the use of walking aides. The innovations for the Ortho3d TKR lie in the metallic components.

Unlike current TKR technologies, the Ortho3d TKR is tailored to mimic the biomechanics of the individual patient. The femoral component and the stemmed tibial plate, shown in Figure 10.1 (a), are manufactured out of nitinol. This offers numerous advantages over current materials. First, nitinol possesses a low stiffness, much closer to that of bone than stainless steel, titanium or cobalt chrome. Second, nitinol is a superelastic alloy. This means that it undergo deformations greater than 5% and completely recover its original shape. This is crucial in the human body because bone also has high elastic properties, capable of recovering up to 2.5% deformation. Stainless steel, titanium, and cobalt chrome cannot even recover 1%. The use of nitinol is only the first stage in technological advancements over current technologies. The second stage involves the use of additive manufacturing, or 3D printing, to produce the Ortho3d TKR.

Through the use of 3D printing, the implant can be tailored to an individual’s needs in two ways: Patient-specific stiffness and patient-specific geometry.

The porosity of the metallic nitinol elements in the Ortho3d TKR can be customized porosity to change the level of stiffness of the implant, shown in Figure 10.1

(b). This is done to match the stiffness of the surrounding bone in the patient to create an environment in the involved knee which is biomechanically accurate. The higher the level of porosity, the less stiff the implant becomes. The stiffness of the surrounding bone 122 in the patient can be determined through a bone mineral test, which is conducted using a specially customized X-ray or computed tomography (CT) scan.

Current total knee replacement technologies are mass-designed with little regard for the shape and size of the individual patients’ bones. This leads to brutal surgeries where much of the patients’ bones must be removed and resurfaced to make the implant fit. Due the additive manufacturing process used to produce the Ortho3d TKR, each implant can be customized to exactly fit the specific patient. Through the use of CT data, the size and shape of the patients’ bones will be digitally characterized and each Ortho3d

TKR will be manufactured to fit. The Ortho3d TKR is the only TKR which can offer any of these advantages.

Figure 10.1: (a) Typical total knee replacement [8] and (b) porous nitinol structure offered in the Ortho3d TKR

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This being our only product, we plan to sub-license to a well-established international medic device company with a proper operational agreement. We are planning to first hit the European market following the receipt of the CE mark. Next, we will begin our clinical trials in the US to obtain FDA approval. We intend to capture the total knee replacement market in Europe and the USA which presents us with an enormous opportunity for growth.

10.5 Industry Overview

Total knee replacements (TKR) are the subject of one of the biggest debates in orthopedics due to the growing number of both replacement and revision surgeries [7].

The primary reason for revision surgeries is aseptic loosing of one or more of the components. The main causes of aseptic loosening are excessive wear between articular surfaces, stress shielding of the bone by the prosthesis, and development of a soft tissue at the bone-implant interface. By utilizing better biomaterials for implant components, implant failure can be mitigated by reducing wear debris, improving load transfer, and providing anchorage between the bone and the implant.

The first knee implant, developed in the 1950’s by McKeever, consisted of a single metal component [4]. In the late 1960’s, two design approaches evolved: anatomical and functional. Anatomical models were designed to preserve the anterior cruciate ligament (ACL) and the posterior cruciate ligament (PCL). Functional models, on the other hand, were permitted “nonanatomical joint surface geometries intended to maximize surface area and reduce polyethylene stress [5]”. Milestones in development include the introduction of the tibial peg in 1973 [6], the first mobile bearing knee 124 implant in 1977 [5], and the porous coated anatomical (PCA) knee was in 1979.

Development has continued with the goals of reducing wear, improving kinematics, and mitigating implant loosening.

Current designs of TKRs (Figure 10.1(a)) incorporate ultra-high molecular weight polyethylene (UHMWPE), ceramic and metal components. Typically, the tibial insert and patellar component are made of plastics such as UHMWPE or cross-linked polyethylene

(XLPE), while the femoral component and the tibial tray are made of metals such as titanium alloys, stainless steel or cobalt chromium molybdenum.

As estimate of the number of knee replacements performed in the United States in the last twenty years is shown in Figure 10.2. It is shown that the number of procedures has increased dramatically since 1990 [205]. This is primary due to the ageing population, however, many young Americans are also being treated with TKRs including injured athletes and wounded warriors. According to the Agency for Healthcare Research and Quality, more than 600,000 knee replacements are performed each year in the United

States [8].

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Furthermore, the demand for primary total knee arthroplasties (Figure 10.3) is projected to grow by 673%, resulting in 3.48 million annual procedures, by 2030 [206].

Total knee revisions (Figure 10.4) are projected to double by 2015 and increase by 601% by 2030.

Figure 10.2: Estimated number of total knee replacements performed in the US from 1990-2010 [205]

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Figure 10.3: Projected primary total knee arthoplasties in the US through 2030 [206]

Figure 10.4: Projected total knee revisions in the US through 2030 [206]

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10.6 Market Analysis

10.6.1 Target Customers

The ultimate beneficiaries of the Ortho3d TKR are people with at least one knee that has been severely damaged by arthritis or injury. Mainly due to aging, disability of the knee may be characterized by reduced biological and/or mechanical integrity of the cartilage and surrounding bone. Affected people suffer severe pain in the involved knee.

The patient, though, is actually the secondary customer for our product and not the focus of our sales. Our primary customer will be physicians and surgeons at orthopedic centers.

Ultimately, it is the doctor’s expertise and preference that decide which surgical implant to use.

As such, our focus will be orthopedic surgeons. The results of extensive laboratory tests will be used to show them that choosing the Ortho3d TKR is beneficial for both the patient and the surgeon. The patient can expect a higher quality of life after installation of the Ortho3d TKR compared to other total knee replacement devices. This should be the principal concern for the attending orthopedic surgeon.

The surgeon can expect other, more direct, benefits as well. First, the Ortho3d TKR is designed to simplify the installation procedure and reduce complications. For the surgeon, this means a decreased workload in the operating room, less training and higher probability of a positive patient outcome. Furthermore, the Ortho3d TKR is competitively priced in the total knee replacement market and will not place an added financial burden on the hospital or the patient. Obviously, the decision of what kind of implant to use is

128 complex on many fronts. In order to facilitate sales and proper implantation of the Ortho3d

TKR, our team will employ trained sales representatives.

The first responsibility of our sales team will be to travel to major orthopedic medical clinics, such as the Cleveland Clinic, and reach out orthopedic surgeons and physicians. Our sales team will be highly knowledgeable in all implant technologies, particularly total knee replacements. They will be able to present our product, explain all of the positive features and experimental results, and answer any questions to alleviate concerns.

The second responsibility of our sales team will involve educating orthopedic surgeons on proper implantation procedures for the Ortho3d TKR. In addition to working with the surgeons pre-surgery, our technical sales team will be available to observe the surgery in the operating room. This will allow for reciprocal feedback between orthopedic surgeons and our company.

10.6.2 Market Size

In 2011, the global knee implants market was USD 8.4 billion and is projected the market to grow to USD 15 billion by 2018 [216]. As shown in Section 10.5, total knee replacement technology is still in its infancy. Extreme growth projections indicate that the medical device market is adopting a friendlier and more accepting position on the technology. Combined with the ageing baby boomer population, this presents a fantastic opportunity to enter the market.

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10.6.3 Competition

In the US, the knee implant market is dominated by a small number of companies

(Table 10.1). The list price for porous total knee implants in 2004 was $8,075. In 2008,

Orthopedic Network News estimated that 13% of the implant cost goes to net income.

The rest of the price components are: selling, general and administrative costs (44%), manufacturing (29%), research and development (6%), and tax (6%) [215].

Company Estimated Market Share Market Cap (2004) [213] Zimmer 29.3% $14 B DePuy 23.4% $180 B Stryker Corp. 19.5% $17 B Biomet 11.4% $8 B Smith & Nephew 9.6% $7 B Other 6.8% - Table 10.1: US 2004 knee implant market share, by company

10.6.4 Estimated Market Share and Sales

The Ortho3d TKR is estimated to gain about 1% of the US market share of the total knee replacement market in the 7th year of existence and then acquire about 0.25% each subsequent year. The market share gain happens for various reasons and advantageous the

Ortho3d TKR holds against its competitors like the patient-specific geometry and stiffness, simple installation, and biomechanical likeness.

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The primary customers would be orthopedic surgeons with experience in knee surgery and total knee implantation. Initially, surgeons in the midwest will be targeted given its proximity to the headquarters and major orthopedic centers, such as the Cleveland

Clinic. Later, the Ortho3d TKR will penetrate to the other areas of the United States aiming to capture more market shares. Currently, more than 600,000 patients are operated for total knee replacements in the United States annually which makes U.S. an attractive place for the Ortho3d TKR.

The Ortho3d TKR will be offered for $9000 in the United States. With regards to the cost of device production, more than 600,000 demands per year for the total knees, and the market share of 1% in the 7th year, 1.25% in the 8th year, and 1.5% in the 9th year, we are projecting an annual sale of $54 million, $67.5 million, and $81 million respectively.

The projected sales can be easily met since the annual growth rate for the total knee market is not accounted for the calculations.

10.7 Marketing Plan

10.7.1 Market Entry and Growth Strategy

The market for total knee replacements is relatively young and growing rapidly.

Technology in the US and abroad is overcoming critical boundaries and we are excited to be leading the charge. Our product offers unique advantages over the competition and we plan to emphasize these differences in marketing.

The FDA approval process in the US takes much longer to navigate than Europe’s equivalent: receiving a CE mark. Therefore, we will obtain the CE mark prior to FDA 131 approval. As such, our primary initial market will be in Europe. Introducing our product in

Europe first has many benefits. First, we will be able to help the population with knee injuries more quickly. Second, we will start earning an income that we can use to pay down loans and fund the FDA approval process. The worldwide markets offer sustainable opportunities but gaining FDA approval is critical to our long term success. As an orthopedic implant, our product is classified as a significant risk device by the FDA and will have to go through the full approval process. Because we are utilizing a new material and manufacturing technique, we are not eligible for 510k approval. Prior to any clinical testing, we will have to file or approval from the FDA and IRB. This process will begin immediately after the company is formed. Due to the complexity of the FDA process, as well as the time and monetary commitments, securing funding and negotiating the approval process will be primary commitments. Finally, it will enable us to start building our brand as soon as possible. Having an established brand and ongoing long-term patient studies internationally will be crucial when introducing our product to the US market.

10.7.2 Price

Because of the unique manufacturing technique we use and patient-specificity we can offer, we consider our product a superior quality implant as compared to the competition. Therefore, we intend to price our product on the high side of the market. At

$9000 per unit, our product is slightly more expensive than the average competitor.

However, we are selling a superior product which costs us more money to develop and manufacture. Furthermore, our product will be competitive enough in price that insurance companies will not be able to refuse coverage. 132

10.7.3 Sales Tactics and Service

As discussed earlier, the tactics that govern implant sales are quite unique. While the patient is the ultimate beneficiary of the product, they are not the focus of our sales.

Due to the amount of technical knowledge required to make an educated decision about buying implant technology, the doctor is primary decision maker when it comes to choosing and buying an implant. Therefore, our sales team will target doctors.

We will select and train a group of highly skilled biomedical engineers to build our sales team. In the field, these sales technicians will visit hospitals and clinics to network with doctors. Not only will they be highly knowledgeable about our product but they will be required to become experts on the competition as well. At hospitals, they will observe surgeries in the OR and discuss with doctors the benefits of using our product.

Additionally, after a doctor makes the decision to use our product, our sales technicians will provide on-site support for the doctors and patient, as well as collect and analyze customer feedback.

10.7.4 Distribution

Every one of our products will be hand delivered to the doctor by one of our sales technicians and handled with utmost care from factory to hospital. This way, we can have staff on-site to answer questions and provide assistance from the very beginning.

Additionally, we will be able to avoid miscues with shipping services.

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10.8 Product Development Plan

10.8.1 Intellectual Property

Based on a preliminary patent search, we found no existing patents for porous nitinol total knee replacement implants, stiffness-tailored total knee replacement implants, or patient-specific total knee replacement implants. Therefore, we believe that the opportunity still exists for us to patent our technology and secure the intellectual property.

However, since the patent process is more easily handled by those more familiar with the industry, we plan on hiring outside legal counsel to conduct a more thorough patent search in the United States and in Europe. After this search is complete, we will either file for patents using our existing design or proceed with design revisions that do not infringe on devices found in the search.

10.8.2 Product Design

Our company will follow a design scheme consisting of the following steps: conceptual design, preliminary design, and detailed design.

During the conceptual design stage, we are working on the first generation of concepts for how we will meet the objectives described in the product description. Before concepts are generated, we will interview multiple orthopedic doctors, as well as patients who have received total knee replacements. We will create a detailed record which will contain all of the parameters that will affect the device’s performance. Due to the patient- 134 specific nature of our product, we also have to consider the development of a design algorithm to efficiently characterize the patients’ needs and produce an implant which is fit for them. During the conceptual design stage, we will begin meeting with radiologists to discuss the use of x-ray, CT, and other methods for bone stiffness and geometry characterization.

In the preliminary design stage, we will immediately begin to analyze the design in order to optimize its properties. We will use contract expertise in Finite Element Modeling

(FEM) to determine stresses in material. This will be done for two reason: (1) to ensure proper operation throughout the life of the device and (2) to match the stresses and strains in the prospective implant site. In order to determine these loads, we will use a software package called AbaqusTM. We will also do a limited number of tests on human subjects in order validate our findings on AbaqusTM software package. Because we will not be asking subjects to do movements outside of typical day to day movements, we expect approval from appropriate Investigational Review Board to be straight forward and timely.

After reviewing the results of our FEM study, we can make changes to the shape and size of the device. The preliminary design of the device will be complete when the development team has generated a design that satisfies the performance specification which can be tailored to any size for patient-specificity. The last step of the preliminary design phase will be the generation of an installation procedure for the Ortho3d TKR device. We will also run an in-house bio-skills lab to train the surgeons with our technique.

In the detailed design phase, we will assign tolerances and allowances for our design so we can develop a standard for our device to be manufactured consistently. Due to the patient-specific nature of our design, these tolerances will be relative to anatomical

135 landmarks on the implant site. No process like this currently exists. All of the devices will be manufactured by selective laser melting. We do not anticipate any manufacturing issues which would prevent the Ortho3d TKR from meeting the market demand.

10.9 Manufacturing and Operations Plan

Ortho3d will be the primary manufacturer for the Ortho3d TKR. Ortho3d will locate itself in Toledo, Ohio. We expect that this location will facilitate the hiring of strong candidates as the University of Toledo is consistently graduation biomedical engineers with expertise and interest in the medical device industry. Furthermore, we anticipate that the University will continue to work with us as they have previously demonstrated a commitment to helping local startups, particularly those which were formed based on University research.

Our medical device packaging will be outsourced to an agency similar to GS

Medical Packaging, who provides third party packaging solutions to the medical industry.

Their knowledge and expertise in packaging will accelerate our time to market and help us make informed decisions on optimizing the shelf life of our product. One of the critical functions that will remain in house is quality assurance. This will allow Ortho3d to approve all devices that are released to market, and in the case of defects, easily pinpoint in which stages of the supply chain they occurred. The final stage of the manufacturing plan consists of three independent trial runs to obtain the CE mark and FDA approval. These trial runs validate our ISO 9001 manufacturing methods and facilitate compliance with the medical device regulatory agencies.

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10.10 Management Team

The primary inventors and founders of Ortho3d will make up the initial management team. These members include Jason Walker, who developed the manufacturing process and conceptualized the idea of the Ortho3d TKR. He will be

President and Chief Technical Officer. Dr. Mohommad Elahinia, a professor in mechanical engineering at the University of Toledo, will serve as the Vice President of

Ortho3d. Dr. Elahinia advised all of the research which led to the development of the

Ortho3d TKR vision. Dr. Sonny Ariss is a professor at UT and specializes in new venture creation. He is also a Fellow at the Center for Technological Entrepreneurship and

Innovation. Dr. Ariss will join the company as the Chief Operations Officer.

10.11 Overall Project Schedule

We intend to incorporate our company immediately. Our company will be formed as an S-Corp. The S-corp is limits the number of shareholders in the company to under

100, but we do not see this an obstacle for our company. The S-corp is advantageous over the C-corp in that it avoids double taxation of the shareholders. In the S-Corp, the shareholders can receive profits at the corporate level free of taxation. Any profits are only taxed at the individual level. Furthermore, the simplicity of transferring and selling stocks with an S-corp is favorable compared to LLC’s. Unlike with LLC’s, the sale of corporate stock has virtually no limitations. Within the first two (2) months we plan to begin acquiring office space, hiring key personnel and establishing all major executive officers.

Within six (6) months we plan to hire all fringe personnel.

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We will conduct all of our own marketing. We plan to launch our marketing campaign exactly one (1) year after we incorporate. This will give us time to build a hire and build a marketing department (six (6) months), and time to develop all marketing channels (six (6) months).

Within the first three (3) months, we plan to complete a comprehensive market analysis and theoretical feasibility study. Preliminary studies have already been conducted and indicate highly positive results.

Within the first six (6) months we plan to finalize the first stage design of the

Ortho3d TKR implant and file relevant patent(s). Within one (1) year, we plan to have completed preliminary in-house testing and have developing the second stage design. The second stage design will be done for the initial marketing launch.

Within two (2) years, we will file for FDA clearance and a CE mark. Within four (4) years, we expect to receive our CE mark and within six (6) years our FDA clearance.

We will officially sell our first product internationally in our fourth (4) year. We will make our first domestic sale in our sixth (6) year. We expect midway through our six

(7) year that our company will reach a financial break-even point.

10.12 Critical Risks and Assumptions

A significant risk that Ortho3d faces is proving to surgeons that our device will work. It is anticipated that there will be some skepticism given the radical nature of our product’s innovations. Not only is our product manufactured from a unique material, but it’s manufactured by a unique method. It is critical that we gain surgeons’ support early

138 in our product lifecycle. We are not the first company to introduce a biomedical implant which utilizes nitinol, so we believe that is the smaller of the two hurdles. Convincing surgeons that additive manufacturing for patient-specific implants is a better strategy is expected to be more difficult. However, we will utilize our relationship with the

University of Toledo and the University of Toledo Medical Center to build relationships with surgeons and biomedical engineers whom we believe are prone to early adoption of radical technologies.

Another risk that out company faces is that CE certification and FDA approval is not secured. Nitinol is an approved biomedical material. However, our manufacturing technique is novel. Therefore, we grade the threat of not receiving FDA approval on time as moderate. However, we grade the threat of not receiving FDA approval at all as moderate. Based on historical tendencies, we expect that the CE mark in Europe will be relatively easy. Therefore, our contingency plan in the case of not receiving FDA approval is to operate strictly in the European market.

The final major risk that our company faces is that sales fail to materialize. While this is typically the largest threat facing a new venture, we plan to license our technology to a major corporation that already has an established reputation in the medical device industry. Therefore, we grade this threat as low. Nonetheless, if constant sales cannot be established, we plan to sale the technology outright in or to recoup capital investments.

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10.13 Financial Plan

It is essential for any start-up company to precisely monitor the major financial expenditures. Our detailed studies on the medical device development and marketing enabled us to effectively meet the fund raising required for a balanced budget. Table 10.2 illustrates the financial calculations for our venture including the expenses and the revenues. As it is shown in Table 10.2, the cost required for the marketing and regulatory activities like clinical trials and bench tests are estimated since these costs are crucial for a medical device company to succeed. Overlooking such expenses can lead to a cost overrun by the start-up company. Cash flow is another essential part of the financial prosperity of any company. As a result, we demonstrate our plan for creation and maintenance of a proper level of cash flow.

We are planning to be actively involved in the fundraising events, especially between series A and series B funding which occurs between the first and last quarters of their respective years. It might seem risky to expect such high fundraising; however, the management team is planning to concentrate on the best ways to raise funds. In case the expected funding is not met, the management team would slow the progression of clinical trials and would increase hiring of the sale forces in Europe.

As shown in Table 10.2, we expect that revenue will exceed expenses midway through the sixth (6) year. This point is known as the break-even point. The break-even point is displayed graphically in Figure 10.5. Based on our analysis of the total knee replacement market, we believe that achieving these levels of sales per year in the United States and abroad are quite achievable.

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Expenses and Revenue per Year 60000 Expenses 50000 Revenue

40000

30000

20000 Dollars ($1000) Dollars

10000

0 0 1 2 3 4 5 6 7 8 Years

Figure 10.5: Break-even analysis for Ortho3d

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Table 10.2: Financial calculations for Ortho3d

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References

1. Elahinia, Mohammad H., et al. "Manufacturing and processing of NiTi implants:

A review." Progress in Materials Science 57.5 (2012): 911-946

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